JP2016518529A - Alloy composition - Google Patents

Abstract

Alloy composition. In mass percent units, 24.0-26.0% Cr, 19.0-22.0% Co, 1.00-1.60% Al, 1.00-1.60% Ti, max Up to 1.50% Fe, up to 0.020% C, up to 0.40% Si, up to 0.50% Mn, up to 0.50% Cu, up to 0.50% Mo, up to 0.50% Nb, up to 0.005% B, up to 0.03% Zr, up to 0.10% V, up to 0.02% Mg, up to 0 Up to 25% Ta, up to 0.010% S, up to 0.015% P, up to 0.20% W, up to 0.02% Ca, up to 0.05% N, Zn up to 80 ppm, Sn up to 60 ppm, Pb up to 20 ppm, and balanced Ni and impure Alloy consisting.

Description

  The present invention relates to alloy compositions, and more particularly to alloy compositions based on nickel. The alloy composition of the present invention is particularly suitable for forming products by casting as opposed to wrought processes (ie, machining, forging, or stamping, etc.). The alloys of the present invention are particularly suitable for use in high-pressure steam turbine applications, such as in power plants.

  The trend towards reducing carbon dioxide emissions and improving the efficiency of coal-fired power plants is related to material development that allows operation at steam temperatures above 700 ° C. with pressures up to 80 MPa, or pressures as high as 100 MPa. Brought about a lot of research being done around the world. At these high temperatures and pressures, steel does not have sufficient strength, so material development instead focuses on multiple nickel alloys.

  The development of nickel alloys has generally focused on two main alloy families: solid solution strengthened alloys and precipitation strengthened alloys. Solid solution alloys have a lower strength than the precipitation hardening of their counterparts, which reduces stress rupture strength, but has good ductility and weldability. Precipitation strengthened alloys are used in multiple jet engines.

  Accordingly, the present invention seeks to provide a new alloy composition. In particular, the present invention seeks to provide a castable alloy composition with the required stress rupture strength for power plant operation above 700 ° C and possibly even above 750 ° C.

  Accordingly, the present invention is based on mass%, 24.0-26.0% Cr, 19.0-22.0% Co, 1.00-1.60% Al, 1.00-1. 60% Ti, up to 1.50% Fe, up to 0.020% C, up to 0.40% Si, up to 0.50% Mn, up to 0.50% Cu, Mo up to 0.50%, Nb up to 0.50%, B up to 0.005%, Zr up to 0.03%, V up to 0.10%, Max 0.02% Up to Mg, up to 0.25% Ta, up to 0.010% S, up to 0.015% P, up to 0.20% W, up to 0.02% Ca, up to N up to 0.05%, Zn up to 80 ppm, Sn up to 60 ppm, Pb up to 20 ppm, and Ba Providing Nsu Ni and alloys comprising a plurality of impurities.

  The alloys of the present invention are precipitation strengthened and thus have higher stress rupture strength than previously used solid solution strengthened nickel alloys. Furthermore, the alloy composition is castable.

  The invention will be described more fully by way of example with reference to the accompanying drawings, in which:

2 is a graph showing a comparison of stress rupture strength versus temperature for various steels and Ni alloys to operate at 550-760 ° C. Several nickel alloys included in the graph have been proposed as potential alloys for plant operation above 700 ° C.

It is a SEM micrograph which shows the crack of the cast alloy after heat processing.

FIG. 3 is an equilibrium phase diagram showing a plurality of phases present in the alloy of FIG. 2.

FIG. 3 is an equilibrium phase diagram showing the MX phase, the eta phase, and the M ₂₃ C ₆ phase of the alloy of FIG.

FIG. 3 is an equilibrium phase diagram showing a plurality of phases present in the alloy of the present invention.

Figure 2 shows the MX composition of the alloy according to the present invention.

It is an equilibrium phase diagram showing the relationship between gamma prime% and niobium concentration.

FIG. 3 is an equilibrium phase diagram showing the relationship of Nb MX% to carbon content.

FIG. 2 is an equilibrium phase diagram showing the combined effect of Nb and C on Nb MX content.

FIG. 4 is an equilibrium phase diagram showing MX predictions for various carbon contents with Nb fixed at 0.25%.

FIG. 4 is an equilibrium phase diagram showing MX predictions for various carbon contents with Nb fixed at 0.25%.

FIG. 4 is an equilibrium phase diagram showing MX predictions for various carbon contents with Nb fixed at 0.25%.

FIG. 4 is an equilibrium phase diagram showing MX predictions for various carbon contents with Nb fixed at 0.25%.

FIG. 6 is the result of the calculation of the shell for alloy 740. FIG.

The influence of tantalum on gamma prime (γ ′) mass% is shown.

Figure 3 shows the expected beneficial effect of Mo on gamma prime (γ ') mass%.

2 is a bar graph comparing the tensile strength of the alloys of the present invention at room temperature with other casting candidate Ni alloys for advanced super supercritical (A-USC) applications.

It is a bar graph which compares the yield strength in room temperature of the alloy of this invention with other casting candidate Ni alloy for A-USC use.

It is a figure which compares the elongation at room temperature and the cross-sectional reduction rate of the alloy of this invention with the alloy 263. FIG.

It is the graph which compared the stress rupture strength plotted with respect to the Larson mirror parameter of the alloy (G130 and G130B) of this invention with the well-known data of the some Ni alloy for casting A-USC used as another candidate.

FIG. 6 is a graph of stress rupture strength plotted against Larson mirror parameters for G130 and G130B for different cross-sectional dimensions.

It is a SEM micrograph which shows the G130B casting alloy by which MX crack is not observed at all after heat processing.

  Explore alloy compositions that have high stress rupture strength at temperatures above 700 ° C (eg, high pressure steam turbine casings for power plants, valves, or multiple components of (auxiliary) valves) and can be formed into products by casting. In doing so, the present inventors examined a plurality of Ni alloys shown in Table 1.

  These alloys are strengthened by either a solid solution strengthening mechanism (SS) or a precipitation hardening mechanism (PH). These mechanisms are briefly described below.

  The solid solution strengthened Ni alloy consists of a face centered cubic (fcc) matrix that is strengthened by the addition of solid solution strengtheners such as Mo, Cr, Co, and W, for example. This hardening method depends on the difference in lattice constant between the austenite structure based on Ni and the solute atoms, and the strengthening efficacy is proportional to this difference. The greater the proportion of solute atoms, the greater the effect of curing. However, the amount of alloy additive is limited by the possibility of sigma phase formation and the associated disadvantages.

Precipitation hardening is achieved in Ni alloys by the formation of three types of precipitation: gamma prime (γ ′), gamma prime (γ ′) (or gamma double prime), and grain boundary borides and carbides. . γ ′ is generally given the chemical formula of Ni ₃ (Al, Ti). However, both nickel and aluminum and titanium can be replaced by multiple other elements.

The inventor has found that the alloy 263, Haynes® 282®, and Inconel® 740, all developed as alloys for wrought, have a particularly high stress rupture strength, as shown in FIG. For the most promising candidates. However, alloy 263 suffers from microstructural instability as soon as it is aged due to the precipitation of eta (η) phase (Ni ₃ Ti), which can adversely affect creep strength. Haynes (R) 282 (R) is stronger among the three alloys, but the highest levels of Al and Ti (Al + Ti> 3.5%) and these elements during implantation in the atmosphere The possibility of oxidation of is significantly increased. This can have the potential to result in more oxide inclusion defects in the casting. During the test of casting alloy 282, it was also known that the integrity of the large cross-sectional microstructure was poor. This results in much worse volume cracking than observed with either alloy 263 or Inconel® 740.

  Inconel® 740 was developed as a casting and wrought product for advanced ultra supercritical boiler tubes. In the wrought process, the alloy is cast as a billet and then undergoes mechanical deformation to form a generally desired shape. Next, a final machining step is taken to achieve the desired shape.

  The composition of Inconel® 740 is derived from the aircraft alloy Nimonic 263 and can be age hardened at about 800 ° C. These two composition formulas are shown in Table 2.

  Unfortunately, the original Inconel® 740 composition appears to have tended to bond microcracks in large section tube work and form an unfavorable “G” phase, which is still an eta (η) phase. However, it tends to be less than that of the alloy 263.

  To solve this problem, the alloy composition of Inconel® 740 was changed to create an alloy called 740H, which was announced at EPRI's “6th International Conference on Advanced Material Technology for Fossil Fuel Power Plants” It was published in a paper titled “Optimization of INCONEL (registered trademark) alloy 740 for Advanced Ultra Supercritical Boilers” in October of the year. This paper describes the changes that special metals have made to the Inconel® 740 chemistry and is detailed in Table 3 below. Note the compositional changes of Si, Nb, Al, and Ti.

  The rationale behind each elemental change is summarized as follows.

  The “G” phase is a Si-rich phase, so Si was reduced to the target specification to prevent its formation. The decrease in Nb content limits the tendency of the material to be susceptible to liquefaction cracking in the heat affected zone. The Al: Ti ratio was increased to increase the stability of gamma prime (γ ′). This is the main reason that cracking is prevented due to the aged and more stable matrix. On the other hand, the reduction of Ti can reduce the ability of the material to produce acicular eta (η) phases when used at 725 ° C.

  The inventor has found that when standard wrought alloy 740 or 740H chemistry is used in the casting, cracking occurs after heat treatment, which results in multiple NDE defects and scrap casting.

  Investigations using SEM have found that cracks have a volumetric nature and are thought to expand from multiple intergranular regions of the alloy. In the grain boundary region between these grains, there are a plurality of Nb-rich MX grains that themselves become cracks during the heat treatment. These particles are the reason for the loss of boundary agglomeration, which is the cause of cracks during which stress is induced in the material during heat treatment. Furthermore, after the solid solution treatment, the spaces formed along the plurality of grain boundaries by partial dissolution of Nb MX generated a plurality of crack expansion sites at the plurality of grain boundaries.

  FIG. 2 shows a SEM micrograph of grain boundaries in cast Inconel 740H after homogenization heat treatment at 1100 ° C. As shown in the figure, there are two grain boundary phases. The inventor has discovered that these are Ti-rich MX phase and Nb-rich MX phase. As shown in the SEM micrograph, the Nb-rich MX phase is cracked. This is considered to be caused by a difference in thermal expansion coefficient between the parent phase and the Nb-rich MX phase.

  A series of heat treatment tests were conducted to confirm the stage of precipitate cracking and to verify whether cracking can be avoided by adjusting the rate of temperature rise and fall. These are shown in Table 4.

  Under all conditions in Table 4, a plurality of Nb-rich MX precipitates were present, and almost all cracked.

  The differential scanning calorimeter (DSC) curve diagram shows that the dissolution temperature of Nb-rich MX could be on the order of 1350 ° C. However, the melting point of the alloy is about 1400 ° C. Therefore, once the Nb-rich MX is formed once in the wrought Inconel® 740 and Inconel® 740H compositions after solidification, melting the mass and removing cracks through the MX is almost It is considered impossible.

  Accordingly, the present inventors have found that MX precipitates are present immediately after the liquid solidifies, so that Nb-rich MX particles are removed from the cast alloy Inconel® 740H only by heat treatment, or the rate of temperature rise and fall is increased. It was concluded that it is impossible to prevent cracking by lowering. Through experiments and thermal analysis, the dissolution temperature of Nb-rich MX was found to be 1350 ° C. Since this melting temperature is too close to the melting point of the alloy, this temperature cannot be utilized.

Based on the above, the present inventor examined whether the composition of the alloy could be changed to reduce the formation of Nb MX carbide. For wrought alloys, particle size can be controlled in that recrystallization can be achieved during the wrought process. It is possible to maintain a small particle size by using carbides and other phases to fix the grain boundaries. The inventor has realized that such control of particle size is not achievable in cast products. This is because recrystallization is not promoted by mechanical deformation. Therefore, the reduction of Nb MX and potentially less M ₂₃ C ₆ phase can reduce the grain boundary aggregation and fraction of some γ ′ phases, but crack removal is more important in cast products .

  The inventor used thermodynamic models to calculate the equilibrium distribution of the phases and the separation behavior of the individual elements as a function of temperature.

  A thermodynamic model calculates the equilibrium distribution of phases and the separation behavior of individual elements as a function of temperature when given the elemental concentration of the bulk alloy. The thermodynamic model includes a plurality of mathematical algorithms that are used to determine a plurality of properties of the alloy phase. Several mathematical algorithms use a database containing thermodynamic data for the alloy system of interest. The database contains essential technical data such as production enthalpy, entropy, chemical potential, interaction coefficient, heat capacity, crystal structure, etc. Multiple thermodynamic calculations are based on Gibbs free energy minimization.

FIG. 3 shows an equilibrium phase diagram of Inconel® 740 having a composition as shown in Table 5 (a composition close to the composition of the Inconel® 740 composition formula), and FIG. The equilibrium phase diagram expanded to emphasize the η phase and the M ₂₃ C ₆ phase is shown.

  As shown in the figure, at high temperatures (more likely to achieve equilibrium), the η and MX phases are in equilibrium. Upon cooling from these temperatures, these phases can remain intact or can precipitate after prolonged aging.

  Table 6 shows how the composition of Inconel® 740 was changed by the inventor. New alloys, designed to produce products by casting, have several important chemistry differences that go beyond conventional wrought compositions, thereby making the alloys much more castable.

  Carbon and niobium were reduced to limit Nb carbide formation (MX) after solidification. Thermodynamic calculations were performed on new alloys with low carbon and niobium concentrations. The equilibrium diagram of the new alloy (FIG. 5) shows that the major phases are nearly identical to the standard Inconel® 740 (shown in FIGS. 3 and 4). However, the eta phase, which is brittle, causes cracking and / or impairs creep strength during use, has been removed. Also, as shown below, the alloys of the present invention having a carbon content up to 0.020% and an Nb content up to 0.50% have a plurality of desired properties.

  Further advantages from Nb reduction are as follows. a) Since Nb is an element with a high potential for separation, the alloy is less prone to segregation at large cross sections. This is shown in FIG. 14, which is a shale calculation showing the non-equilibrium separation tendency of the cast alloy 740. b) The reduction of Nb makes the atmospheric nitrogen soluble in the product less soluble, allowing lower nitrogen levels to be achieved than is achieved by standard products without special processes.

  FIG. 5 is an equilibrium phase diagram of the new alloy composition formula 1 in Table 6 and shows that the thermodynamic prediction of the main phase of the alloy is similar to that of alloy 740H. However, the eta phase has been completely removed.

  FIG. 6 uses the same thermodynamic model to predict the composition of the MX phase in the alloys of the present invention. The MX composition of the alloys of the present invention differs from the standard Inconel 740H® composition. The alloy of the present invention has an Nb mass fraction increase of less than 0.01 at temperatures above 900 ° C., and the Ti mass fraction is substantially constant from 300 ° C. to the melting point. Thus, the predicted total mass fraction of MX is significantly reduced.

  Since the eta phase is detrimental, it is further reduced in cast alloys compared to 740H due to Nb reduction.

Due to the decrease in Nb concentration, the gamma prime (γ ′) fraction is lowered in the alloy of the present invention as compared with the case where the Nb content is higher, such as a conventional wrought alloy. Nb substitutes for Al in (γ ′) (Ni ₃ Al), just like Ti. The wrought alloys 740 and 740H contain about 21% gamma prime (γ ′). At this level of precipitation, along with the high level of MX precipitation, the inventors have found that the cast alloy has very low room temperature ductility. This is an important factor to consider when the plant where the alloy is used repeats thermally between shutdown and start-up and when repairs are required during use, for example by welding. is there. Thus, in the alloys of the present invention, gamma prime (γ ′) is preferably increased to a level where ductility is balanced with the ultimate creep strength.

  In a preferred embodiment, the total mass% of aluminum and titanium ensures that 2.8% by mass, preferably greater than or equal to 2.9% by mass, and / or ensures that the total amount of Ta is 0. .13 to 0.17% by weight and / or having an amount of Nb present in an amount of 0.23 to 0.27% by weight and / or reducing the amount of Mo to 0.23 By increasing to ˜0.27% by weight, the amount of gamma prime (γ ′) is increased.

  The alloys of the present invention are, in mass percent units, 24.0-26.0% Cr, 19.0-22.0% Co, 1.00-1.60% Al, 1.00-1. 60% Ti, up to 1.50% Fe, up to 0.020% C, up to 0.40% Si, up to 0.50% Mn, up to 0.50% Cu, Mo up to 0.50%, Nb up to 0.50%, B up to 0.005%, Zr up to 0.03%, V up to 0.10%, Max 0.02% Up to Mg, up to 0.25% Ta, up to 0.010% S, up to 0.015% P, up to 0.20% W, up to 0.02% Ca, up to N up to 0.05%, Zn up to 80 ppm, Sn up to 60 ppm, Pb up to 20 ppm, and rose Consisting scan Ni and alloys comprising a plurality of traces additives and more impurities.

  The term “consisting of” is used herein to indicate that 100% composition is mentioned and the presence of additional components is excluded. As a result, the sum is 100%.

  The rationality of the composition of the present invention is demonstrated below. All amounts are weight percent unless otherwise stated.

  Nickel: Balance.

  This is the matrix element that forms the balance of the chemistry of the material and has good solubility in a wide variety of secondary elements and as a result can be easily alloyed. Face centered cubic (FCC) nickel lattices are compatible with both solid solution and precipitation strengthening mechanisms.

  Chromium: 24.0-26.0 mass%, preferably 24.5-25.0 mass%.

Chromium provides the alloy with extremely important corrosion resistance by rapidly forming a surface layer (Cr ₂ O ₃ ) that effectively protects the underlying alloy from excessive corrosion attack under high temperature corrosion. . Chromium also promotes the formation of a surface alumina layer that further enhances the corrosion resistance of the alloy at low temperatures. Together with Mg (up to 0.02%) and Si (up to 0.4%), the protective properties of the scale are further enhanced and become effective against higher temperatures. The function of these trace elements is to increase scale adhesion, scale density, and scale degradation resistance.

  Chromium has a mild solid solution strengthening effect on the parent phase due to the mismatch of atomic parameters with Ni.

  The minimum level of Cr is selected to ensure the formation of α-chromia scale at high temperatures. This Cr level is about 24.0%, preferably 24.5%. Slightly higher Cr levels accelerate α-chromia formation, but do not change the nature of the scale. The maximum level of Cr is determined by the stability of the alloy. This maximum level of Cr was found to be about 26.0%, or better 25.0%.

  Cobalt: 19.0 to 22.0 mass%, preferably 19.5 to 20.5 mass%.

The main purpose of the Co addition is to indirectly increase the solidus temperature by improving the solubility of multiple γ ^' precipitates, followed by increasing the working temperature of the alloy. This increase increases the γ ^' volume fraction of the alloy and has a further strengthening effect, so an amount greater than 19.0%, preferably greater than 19.5% is used.

  The atomic size of cobalt, similar to the atomic size of nickel, limits the solid solution strengthening effect of cobalt on the matrix and, due to cost, the amount is 22.0% or less, preferably 20.5% Or less.

  Aluminum: 1.00 to 1.60% by mass, preferably 1.20 to 1.55% by mass, more preferably 1.40 to 1.55% by mass.

Aluminum is a major alloying element that partitions a plurality of γ ^' precipitates. ^'As forming element, Al is ^gamma' main gamma volume fraction, and as a result, because it has a direct effect on the precipitation strengthening of the alloy, Al is at least 1.00 wt%, preferably at least 1.20 mass %, More desirably in an amount of at least 1.40% by weight.

Aluminum provides corrosion resistance to the alloy by forming an alumina (Al ₂ O ₃ ) surface protection layer under high temperature corrosion. Since an excessive amount of aluminum can reduce the ductility of the alloy, the amount is limited to 1.60% by weight or less, desirably 1.55% by weight or less.

  Titanium: 1.00 to 1.60% by mass, preferably 1.20 to 1.55% by mass, more preferably 1.40 to 1.55% by mass.

Titanium ^'are forming element, ^gamma' another major gamma so also tend to distribute the phase, Ti is at least 1.00%, preferably at least 1.20%, more preferably in an amount of at least 1.40% Exists.

  Ti in an amount exceeding 1.60% is liable to cause internal oxidation that causes a decrease in the ductility of the parent phase. Preferably Ti is present in less than 1.55%.

  Titanium shares a precipitation strengthening effect similar to aluminum.

  In order to further compensate for the decrease in the gamma prime (γ ′) fraction due to the Nb content decrease, the mass percentage of Al and Ti is increased in the new alloy. Table 13 shows the effect of increasing Al + Ti mass% on the estimated gamma prime (γ ′) fraction, as well as balance chemistry, as detailed in the proposed new composition formula. G130 is as shown in Table 14. Desirably, Al + Ti mass% is greater than or equal to 2.8 mass%, more desirably greater than or equal to 2.9 mass%, or even greater than or equal to 3.0 mass%. In order to reduce the possibility of oxidation during implantation in the atmosphere, the Al + Ti mass% is preferably kept below 3.15 mass%, or more preferably below 3.05 mass%.

  Niobium: 0.000 to 0.50 mass%, preferably 0.20 to 0.30 mass%, and more preferably 0.23 to 0.27 mass%.

Niobium promotes the formation of the γ ^' phase, but to a lesser extent compared to Al and Ti. The present invention reduces the Nb level in the alloy by 0.50% or less, or less than 0.50%, or preferably less than 0.40%. The presence of Nb leads to an increase in the volume fraction of γ ^′ , so a level of 0.15% or greater, more desirably a level of at least 0.20% by mass, or even at least 0.23% by mass is desirable. Can be.

The main advantage of Nb reduction is to prevent the precipitation of Nb rich MX (NbC). Another advantage is to reduce the tendency of eta phase formation. In addition, Nb has a high separation tendency in a large cross-section casting, and limiting the concentration helps to reduce this tendency. In the present invention, the amount of Nb up to 0.30% by mass in the gamma prime (γ ^′ ) volume fraction is obtained without forming Nb-rich MX or without having a harmful tendency to eta phase precipitation during use. This produces the desired slight increase. However, the maximum amount of Nb up to 0.49%, or up to 0.40%, more preferably up to 0.30%, most preferably up to 0.27% by weight, keeps any possibility of MX precipitation low. Is preferred.

  Carbon: 0.000 to 0.020 mass%, preferably 0.005 to 0.015 mass%.

Carbon is an alloy element that distributes a plurality of grain boundary regions. Its primary function is to form various carbides with other alloying elements to provide increased grain boundary aggregation. Predicted carbides in the alloys of the present invention are M ₂₃ C ₆ and MX. Moreover, when carbon is dispersed in the matrix, it has an appropriate solid solution strengthening.

  The upper limit of 0.020 wt% limits the formation of Nb-rich MX, especially when niobium is also reduced. The inventor has found that carbon levels above 0.02% will result in unacceptable levels of Nb MX in the casting, particularly in cross-sectional dimensions above 100 mm. Preferably, the carbon is in an amount of at least 0.005% by weight to achieve a plurality of grain boundary carbides, thereby increasing grain boundary aggregation and providing a certain solid solution strengthening when dispersed in the parent phase. Exists. Preferably, the carbon is present in an amount of 0.015 wt% or less to avoid excessive carbide precipitation at the grain boundaries.

  Boron: 0.000 to 0.005 mass%, zirconium: 0.000 to 0.03 mass%.

  Boron and Zr can be added to suppress carbide coarsening and consequently increase grain boundary strength. Creep strength and ductility increase with the addition of these elements. The addition of boron improves the grain boundary slide resistance.

  Manganese: 0.00 to 0.50 mass%, preferably 0.15 to 0.45 mass%.

  Manganese improves the weldability of the alloy by reducing the solidification range of the alloy. Therefore, Mn is present up to 0.50% by weight, preferably at least 0.15% by weight. However, manganese is preferably present in a maximum amount of 0.45 wt% or less because it reduces the integrity of the protective scale.

  Silicon: 0.00-0.40 mass%, preferably 0.10-0.25 mass%.

  Silicon controls the carbide morphology and may be beneficial for corrosion resistance. If the Si level is too high, this can then promote G phase formation.

  Furthermore, an excessive amount of Si can contribute to loss of ductility, toughness, and workability. Accordingly, Si is preferably present up to a maximum of 0.40% by weight to control morphology and carbide content and limit loss of ductility, toughness, and workability. A preferable range is 0.10 to 0.25% by mass.

  Molybdenum: 0.00 to 0.50 mass%, preferably 0.20 to 0.30 mass%, more preferably 0.23 to 0.27 mass%.

The present invention limits the Mo content to 0.50% due to its tendency to separate and the ability to form Mo-rich precipitates (carbides). Mo contributes to solid solution strengthening of the parent phase. Mo is mainly distributed within the gamma (γ) matrix, M ₂₃ C ₆ , and gamma prime. Mo contributes to solid solution strengthening of the gamma matrix and a slight increase in gamma prime, and therefore preferably in the range of 0.20 to 0.30% by mass, more preferably 0.23 to 0.28% by mass, ideal Specifically, it is added at 0.25% by mass, and compensation is performed to overcome the strength loss due to the removal of Nb. FIG. 16 shows the expected small but beneficial effect of Mo on increasing gamma prime (γ ′) mass%.

  Since casting cannot be homogenized to the same extent as a wrought alloy, further restrictions on the elements that separate into the cast alloy are more important compared to the wrought alloy.

  Nitrogen: 0.00-0.05% by mass, preferably at most 0.03% by mass.

Titanium is the most potent nitride-forming element, so it easily combines with nitrogen to form MX. Keeping residual nitrogen to a minimum prevents the formation of unwanted MX, which has a diminishing effect on the gamma prime γ ^′ mass fraction.

  Sulfur and phosphorus are unnecessary trace elements and are distributed to a plurality of grain boundaries of the Ni alloy. Therefore, these elements are desirably controlled to a low level so as not to affect the grain boundary strength.

  Sulfur is desirably controlled up to 0.010% by weight, preferably up to 0.005%, ideally less than or equal to 0.001% by weight and phosphorus up to 0.015% by weight, preferably Is controlled to a maximum of 0.010%.

  Zinc, tin, and lead are low melting point constituents and their presence can affect grain boundary aggregation and are therefore considered to be trump elements in the Ni alloy. Desirably, these elements are limited to 80 ppm Zn, 60 ppm Sn, and 20 ppm Pb, and more desirably 30 ppm Zn, 50 ppm Sn, and 10 ppm Pb.

Tantalum is a carbide forming element in Ni alloys and is desirably limited to a maximum of 0.25%. The inventor has realized that, according to the simulations performed, tantalum increases the gamma prime (γ ^′ ) mass fraction by 0.2% with the addition of 0.15 wt%. It can also be seen from multiple equilibrium calculations that Ta only distributes in gamma (γ) and gamma prime (γ ^′ ). Therefore, the addition of Ta will not affect the unwanted increase in MX. Therefore, in the alloy of the present invention, Ta is preferably present at a level of 0.10 to 0.20 wt%, more desirably 0.13 to 0.17 wt%. FIG. 15 shows the effect of tantalum on gamma prime (γ ′) mass%.

  Copper is considered a residue, not the intended additive of the present invention, and is therefore limited to 0.50 wt%, preferably 0.10 wt%.

  Iron has some solid solution strengthening properties, but in the present invention it is limited to less than 1.50% by weight, preferably less than 1.00%, more preferably less than 0.80% and is considered a residue.

  Vanadium and tungsten are not intentionally added in the present invention and are limited to the following values. W is at most 0.20% by mass, preferably at most 0.05% by mass. Vanadium is a maximum of 0.10% by mass, desirably a maximum of 0.05% by mass.

  Magnesium can be used as an oxygen scavenger or desulfurizer. Mg levels can be allowed up to 0.02% by weight in the alloy. However, cast alloys should preferably be limited to 0.01% by weight due to the tendency to produce oxide films when cast in air.

  Calcium can be added as a desulfurizer and oxygen scavenger. In this respect, only a maximum of 0.02% by weight, desirably a maximum of 0.01% by weight, is acceptable.

  A list of the main phases present in the cast alloy being characterized.

  Mother phase: gamma (γ).

  The alloy matrix consists of γ (gamma) with γ ′ up to 20%.

  Gamma prime (γ ').

  The γ 'phase provides strong precipitation strengthening for the alloy.

  MX

  Decreasing the Nb content and reducing the carbon content below or even below 0.50% means that no Nb-rich MX is formed. Thus, the MX mass fraction is lower than for alloys 740 and 740H.

M ₂₃ C ₆ is a Cr-rich precipitate that typically occurs along grain boundaries, as in many other nickel-based superalloys. It precipitates as a function of lower temperature aging.

  Sigma: Conventionally, this phase has a high Cr content and is a low temperature phase.

  Sigma was not detected under the solid solution treatment conditions and aging in the present invention.

M ₃ B ₂ : not detected.

  In the casting industry, the successful production of solid solution strengthened alloys as large section castings is currently commercialized.

In gamma prime (γ ^′ ) reinforced alloys, these alloys are relatively new, so there is little global experience in casting these new materials and no casting specifications are available.

  Special considerations are desirably given below to these alloys.

  a) Melting technique: The melting technique of Ni alloys differs in many respects from conventional steel production. Special techniques can be used to prevent oxidation of multiple volatile alloy elements. VOD, VODC, or AOD melting is used for the manufacture of multiple large components.

  b) Injection technique: Ni alloys alloyed with titanium and aluminum and injected in the atmosphere are preferably injected with care to prevent multiple oxidation defects on the surface of the casting.

  c) Solidification: The solidification characteristics of Ni alloys are quite different from steel, and conventional process techniques are often not satisfactory in producing simply acceptable quality castings. A process for casting nickel alloys that can be used with the alloys of the present invention is disclosed in WO 2009/130472.

  d) Heat treatment: The alloys of the present invention are ideally heat treated by a solid solution treatment followed by water quenching, air cooling, or cooling with static air cooling if the cross-sectional dimensions are acceptable, or for specific applications Is "as cast" without heat treatment. Special homogenization techniques can be used for specific cross-sectional dimensions.

  Accordingly, the present invention provides a new alloy composition for precipitation hardening superalloys that has high stress rupture strength at temperatures above 700 ° C. and can be cast as a casting. Thus, the alloys of the present invention are desirable for the production of multiple steam turbine casings, valves, and (auxiliary) valve components using a casting process, or for high temperatures, strength, and corrosion resistance above 700 ° C. Ideal for use for any other application to be done. In one embodiment, an alloy is used in a valve or valve component for operation at 500-850 ° C, preferably 650-750 ° C. The (auxiliary) valve component may be a component such as a bonnet, cover, bend, or elbow and may operate at high temperatures within the power plant. Tables 7 and 8 below summarize the present invention.

  Trump element: (maximum value)

  Zn: up to 80 ppm, desirably up to 30 ppm, Sn: 60 ppm, desirably up to 50 ppm, Pb: up to 20 ppm, desirably up to 10 ppm

  The present invention is particularly applicable to one-off castings, such as those typically made using sand molding casting techniques in which resin-bonded particulate material is used to create the mold. . The sand-type particulate material may be, for example, sand, zircon, fused quartz, alumina, magnesia, ceramic spheres, or chromite, or any combination thereof. For example, the alloy can be used to produce high pressure steam turbine casings, or valves, for example used in power plants, by casting. In the latest advanced ultra-supercritical turbines, the operating temperature of components exceeds 700 ° C.

  A series of thermodynamic predictions are made, and there are compositional formulas for all major elements, as shown in Table 9. The results are in FIGS. 7 to 13 and Tables 10 to 12.

  FIG. 7 shows changes in the maximum γ ′ content from 400 to 1600 ° C. on the x-axis and changes in the Nb content on the y-axis at a fixed carbon content of 0.015%. Inconel® 740 and 740H (assuming 0.015% carbon) have a γ ′ content of 21% and 19.7%, respectively, while the present invention with an Nb content of 0.5% It has a γ ′ content of about 18%, and an Nb content of 0.25% has a γ ′ content of about 17%. Thus, the alloys of the present invention have good multiple precipitation hardening properties.

  FIG. 8 shows the change in maximum mass% of the Nb MX phase between 800 and 1350 ° C. on the y-axis and the change in mass% of carbon on the x-axis at a fixed Nb content of 0.25%. As shown in the figure, the mass fraction of Nb MX is 0 at a carbon content of less than 0.015%. This is why the preferred upper C level is 0.015% or less. However, even at carbon levels up to 0.02%, Nb MX is present in less than 0.02% by weight. FIG. 9 is similar to FIG. However, FIG. 9 shows the combined effect of Nb and C content by plotting the Nb content on the x-axis at the top of the graph and the C content on the x-axis at the bottom of the graph. There is a large amount of Nb MX on the left side, but as both Nb and C are reduced, the level of Nb MX decreases. At Nb content of 0.43% and less, and C content of 0.015% and less, it is expected that very little or no Nb MX will be produced.

  The thermodynamic calculations show that:

  1) The eta phase is completely removed from the alloys of the present invention by reducing the carbon to 0.018% or lower and by reducing Nb to 0.54 or lower.

  2) In the present invention, Nb MX is completely removed when carbon is 0.015% or lower and Nb is 0.43% or lower, and gamma prime is expected to be 17%.

  10-13 show the effect of reducing the carbon content on Nb MX and Ti MX. All graphs assume an Nb content of 0.25%, and the carbon content is 0.046%, 0.036%, 0.029%, and FIG. 10, 11, 12, and 13, respectively. 0.015%. As shown in these figures, the amount of Nb in the MX phase (and the amount of Nb MX as part of the overall MX) decreases as the Nb content is reduced.

  Both Nb and C are important in determining whether NbC, which has been shown to crack in large section castings, will occur. Therefore, as in the present invention, it is beneficial to keep the Nb and C contents low.

  Tables 10-12 show the results of thermodynamic calculations based on the compositions of Table 9 (similar to FIGS. 7-13) with varying Nb and C contents. These tables show the maximum amount of Ti and Nb in any MX phase and the maximum amount of eta phase in the temperature range of 800-1350 ° C. Examples with * are outside the scope of the present invention.

  Therefore, since C has the greatest influence on the compositional distribution of MX, desirably less C is selected.

  It can be seen that the eta phase appears when Nb increases from 0.62% by weight, and that the percentage of gamma prime increases from 17% to 21% when Nb increases from 0.25% to 1.86% by weight. .

  Therefore, Nb has the greatest influence on the gamma prime and eta phase.

  The eta phase disappears completely when the composition of C is 0.018% and lower and the composition of Nb is 0.54% and lower. The Ti and Nb distribution lines are nearly flat at a composition of 0.015% and lower C, and 0.43% and lower Nb. During the calculation, the gamma prime fraction decreases from 20% to 17% as the Nb content decreases.

  7 to 13 and Tables 10 to 12 show that C has an influence on MX that is greater than Nb. As C increases, the predicted NbC increases. Nb also affects the eta phase and gamma prime, and the ratio of eta phase and gamma prime increases with increasing Nb content. Thus, maintaining C at a low content (depending on practicality) while selecting an appropriate Nb content provides sufficient gamma prime and no eta phase.

  From FIGS. 7-13 and Tables 10-12, the maximum amount of carbon of 0.02% or lower and the maximum amount of Nb of 0.5% or lower removes the eta phase and greatly reduces the amount of Nb MX. Can be found to maintain a low level. At Nb content of 0.25% or lower, or C content of 0.015% or lower, the presence of eta and Nb MX is eliminated.

  The inventor has realized that a slight increase in the fraction of γ ′ can increase the strength of the alloy, which can now be achieved without adversely affecting ductility due to a decrease in MX precipitation. In order to realize an increase in the fraction of γ ′, the present inventors investigated the influence of the total amount of Ti and Al in the alloy. Table 13 shows the results of equilibrium simulation at room temperature. In it, the amount of Ti and Al can be varied, otherwise the composition is equal to alloy G130 in Table 14.

  It can be seen that by increasing the Al content to 1.43 wt% and the Ti content to 1.52 wt%, the gamma prime (γ ′) can be increased by 10%.

  According to the results in Table 13, in order to prevent precipitation of the eta (η) phase during long-term use aging, Al + Ti is 2.8% by mass or more, more preferably 2% while maintaining Ti at a maximum of 1.6% by mass. .9% by mass or more is desirable.

To further assist in compensating for the decrease in the predicted gamma prime (γ ^′ ) fraction in the new alloy, 0.15 wt% Ta is added at an acceptable level up to 0.25 wt%. This has the effect of increasing the gamma prime (γ ^′ ) mass%, as shown in FIG. Here, gamma prime (γ ^′ ) mass% increases Ta (mass%) while maintaining Al, Ti, Nb, and Mo in the proposed composition formula 1 as detailed in Table 6. Plotted against. The gamma prime (γ ^′ ) fraction increases by 0.2% with the addition of 0.15 wt% Ta. From the calculation, Ta only distributes in gamma (γ) and gamma prime (γ ^′ ). Therefore, the addition of Ta has no effect on the increase in MX.

  Table 14 shows the composition formulas of three test alloys included in the scope of the present invention in mass%. The rightmost column shows the estimated mass fraction γ ′ of each alloy.

  A G130 melt was first made and tested. The main goal is to ensure the microstructural integrity of the material from thin to thick sections, as may be required for high quality turbine components. Alloy G130 shows that reducing the carbon content from alloy 740H and removing Nb reduces MX precipitation to such an extent that intergranular cracks and volumetric defects are removed.

The G130 alloy has lower strength than G130A and G130B, which is believed to have been caused by the lower γ ^′ mass fraction. G130A and G130B address this without adding most of MX to the alloy matrix (especially in the large cross section).

The increase in the γ ′ fraction over G130 is due to the introduction of nominally 0.25% by weight of Nb (G130A composition) and the total content of Al + Ti further increased to 2.8% by weight and nominally 0.8. This was realized by adding 15% by mass of Ta (G130B composition). 7 ^'effect of Nb on the mass fraction, Table 13 ^gamma' gamma effect of Al + Ti into the mass fraction, FIG. 15 shows the effect of Ta on the gamma ^'mass fraction. During all tests, Mo was the desired additive and was added to a level of 0.25% by weight. FIG. 16 shows the effect of Mo on the γ ^′ mass fraction.

  Table 15 summarizes the main differences in composition and properties between G130, G130A, and G130B.

  Tables 16 and 17 detail the room temperature and high temperature mechanical properties of alloys G130, G130A, and G130B having cross sections with thicknesses of 100 mm and 200 mm. The UTS for tensile test, yield strength, elongation, cross-sectional reduction rate, and Charpy test results in the case of room temperature test are given.

  As shown, the mechanical properties of G130B are superior to the mechanical properties of G130A which are superior to the mechanical properties of G130.

  FIG. 17 illustrates the advantages of thick section tensile strength of G130B, G130A, and G130 alloys compared to other precipitation hardened Ni alloys manufactured as castings. The alloy of the present invention has an improved UTS in the 200 mm cross section tested compared to the cast alloys 282, 263 and 740H. On the other hand, the tested alloys of the present invention are comparable to alloys 282, 263 and 740H in the 100 mm cross section. However, with the 263 and 282 alloys, their stretch and cross-section reduction properties are much poorer, making it very difficult to process, weld repair and manufacture these materials. This can mean that these materials have certain maximum cross-sectional dimension limits.

  FIG. 18 details the yield strength at room temperature of the alloys of the present invention. The yield strength is comparable to other A-USC casting alloy candidates.

  Alloy 263 is known as a candidate for 750 ° C. A-USC application, although it may have the disadvantage of instability during long-term use. Such materials are generally known for generally better ductility and better weldability characteristics than other precipitation hardened Ni alloys. However, FIG. 19 shows that the ductility of the large section of the material is inferior to that of the G130B alloy when cast as a casting. On the other hand, as FIG. 18 shows, the strength of the alloy of the present invention at room temperature is comparable to or higher than that of the alloy 263.

  FIG. 20 shows the stress rupture performance of alloys G130 and G130B after a short term test compared to other cast Ni alloys. This indicates that alloy G130 performs better than solid solution strengthened alloys, but, as expected, has lower strength than other precipitation strengthened alloys. However, alloy G130 is suitable for use because it does not suffer from microcracking at large cross-sections that would be problematic for alloys 263, 740, and 282 and prevent their use as large cross-section castings. Also, G130 does not have the disadvantage of forming an undesired eta phase during use, which is a problem with alloys 740 and 263 of standard composition. The results for G130B (preferred alloy range) in FIG. 20 show higher stress rupture strength than G130, and the stress test results are close to the scatter band of the cast alloy 263. However, it should be noted that the data for cast alloys 263, 740, and 282 are taken from a very small 7 kg cast and do not represent the stress rupture strength in the original thick section component. . When 263, 740, and 282 are tested with larger specimen sizes, eg, castings over 500 kg, most stress rupture samples are premature due to very poor microstructure grain boundary integrity. Or even broken during loading.

  However, the G130 and G130B results are from a suitable large test material and thus relate to the original component properties. Furthermore, no premature failure was observed.

  FIG. 21 details short-term stress rupture versus Larson mirror parameters for different cross-sectional dimensions of two alloys (G130 and G130B) within the scope of the present invention. The graph shows that the A-USC target criteria of 80 Mpa at 100,000 ° C. and 100,000 hours can be clearly filled with all the materials tested, and that the G130B melt whose sample is still under test (preferred composition) ) Indicates that there appears to be a possibility of achieving 80 MPa in applications at 750 ° C. and 100,000 hours.

  FIG. 22 is a SEM micrograph of the grain boundary of alloy G130B after casting and heat treatment, showing that no MX carbide precipitates are present.

Claims (33)

  In mass percent units, 24.0-26.0% Cr, 19.0-22.0% Co, 1.00-1.60% Al, 1.00-1.60% Ti, max Up to 1.50% Fe, up to 0.020% C, up to 0.40% Si, up to 0.50% Mn, up to 0.50% Cu, up to 0.50% Mo, up to 0.50% Nb, up to 0.005% B, up to 0.03% Zr, up to 0.10% V, up to 0.02% Mg, up to 0 Up to 25% Ta, up to 0.010% S, up to 0.015% P, up to 0.20% W, up to 0.02% Ca, up to 0.05% N, Zn up to 80 ppm, Sn up to 60 ppm, Pb up to 20 ppm, and balanced Ni and multiple Alloy consisting of impurities.   2. An alloy according to claim 1 consisting of less than 0.50% Nb, or less than 0.49% Nb, preferably less than 0.40%.   3. An alloy according to claim 1 or 2, comprising up to 0.25% Nb, preferably up to 0.15% Nb.   2. Alloy according to claim 1, consisting of 0.20 to 0.30% Nb, preferably 0.23 to 0.27% Nb.   5. An alloy according to any one of the preceding claims, consisting of up to 0.015% C.   6. An alloy according to any one of the preceding claims, consisting of 0.10 to 0.20% Ta, preferably 0.13 to 0.17% Ta.   7. An alloy according to any one of claims 1 to 6 consisting of 0.005% or more C.   8. An alloy according to any one of the preceding claims, comprising at least 2.8% by weight Ti + Al, desirably at least 2.9% by weight Ti + Al.   9. An alloy according to any one of the preceding claims, comprising less than 3.15% by weight Ti + Al, desirably less than 3.05% by weight Ti + Al.   10. An alloy according to any one of claims 1 to 9 consisting of 1.40 to 1.55% Al.   11. An alloy according to any one of the preceding claims, consisting of 1.40 to 1.55% Ti.   12. An alloy according to any one of the preceding claims, consisting of 0.20 to 0.30% Mo, preferably 0.23 to 0.27% Mo.   13. An alloy according to any one of the preceding claims, consisting of 0.10 to 0.25% Si.   14. An alloy according to any one of the preceding claims, consisting of 0.15 to 0.45% Mn.   15. An alloy according to any one of the preceding claims consisting of 24.5 to 25.0% Cr.   16. An alloy according to any one of the preceding claims, consisting of 19.5 to 20.5% Co.   17. An alloy according to any one of the preceding claims, consisting of up to 0.10% Cu.   18. An alloy according to any one of the preceding claims, consisting of up to 1.00% Fe, preferably up to 0.80% Fe.   The alloy according to any one of claims 1 to 18, comprising Mg.   20. An alloy according to any one of the preceding claims, consisting of up to 0.03% N.   21. An alloy according to any one of the preceding claims, comprising up to 0.005% S, more preferably 0.001% S.   2. An alloy according to any one of claims 1 to 21 consisting of up to 0.010% P.   23. An alloy according to any one of the preceding claims, consisting of up to 0.01% beloved Ca.   24. The alloy of any one of claims 1 to 23, comprising up to 0.05% W.   25. An alloy according to any one of claims 1 to 24, comprising up to 0.05% V.   26. Alloy according to any one of the preceding claims, consisting of up to 30 ppm Zn, up to 50 ppm Sn, up to 10 ppm Pb.   Cast product comprising an alloy according to any one of claims 1 to 26.   A steam turbine casing or a gas turbine casing or valve comprising the alloy according to any one of claims 1 to 26.   A valve or valve component comprising the alloy of any one of claims 1 to 26.   30. A steam turbine casing or gas turbine casing or valve or valve component according to claim 28 or 29, wherein the component is a cast product.   31. The valve or valve component according to claim 30, wherein the valve or valve component operates at 500-850C, desirably 650-750C.   An alloy substantially as hereinbefore described and illustrated with reference to the accompanying drawings.   A valve, valve component, steam turbine casing or gas turbine casing, substantially as hereinbefore described and illustrated with reference to the accompanying drawings.