Patent Application
20160312341
The invention relates to a nickel-chromium-titanium-aluminum wrought alloy with very good wear resistance and at the same time very good high-temperature corrosion resistance, good creep strength and good processability.
Austenitic age-hardening nickel-chromium-titanium-aluminum alloys with different nickel, chromium titanium and aluminum contents have long been used for outlet valves of engines. For this service, a good wear resistance, a good high-temperature strength/creep strength, a good fatigue strength and a good high-temperature corrosion resistance (especially in exhaust gases) are necessary.
For outlet valves, DIN EN 10090 specifies especially the austenitic alloys, among which the nickel alloys 2.4955 and 2.4952 (NiCr20TiAl) have the highest high-temperature strengths and creep rupture stresses of all alloys mentioned in that standard. Table 1 shows the composition of the nickel alloys mentioned in DIN EN 10090, while Tables 2 to 4 show the tensile strengths, the 0.2% offset yield strength and reference values for the creep rupture stress after 1000 h.
Two alloys with high nickel content are mentioned in DIN EN 10090:
Compared with NiFe25Cr20NbTi, NiCr20TiAl has significantly higher tensile strengths, 0.2% offset yield strengths and creep rupture stresses at higher temperatures.
EP 0639654 A2 discloses an iron-nickel-chromium alloy consisting (in weight-%) of up to 0.15% C, up to 1.0% Si, up to 3.0% Mn, 30 to 49% Ni, 10 to 18% Cr, 1.6 to 3.0% Al, one or more elements from Group IVa to Va with a total content of 1.5 to 8.0%, the rest Fe and unavoidable impurities, wherein Al is an indispensable additive element and one or more elements from the already mentioned Group IVa to Va must satisfy the following formula in atomic-%:
0.45≦Al/(Al+Ti+Zr+Hf+V+Nb+Ta)≦0.75
WO 2008/007190 A2 discloses a wear-resistant alloy consisting (in weight-%) of 0.15 to 0.35% C, up to 1.0% Si, up to 1.0% Mn, >25 to <40% Ni, 15 to 25% Cr, up to 0.5% Mo, up to 0.5% W, >1.6 to 3.5% Al, >1.1% to 3% in the total of Nb+Ta, up to 0.015% B, the rest Fe and unavoidable impurities, wherein Mo+0.5 W is ≦0.75%; Ti+Nb is ≧4.5% and 13≦(Ti+Nb)/C≦50. The alloy is particularly useful for the manufacture of outlet valves for internal-combustion engines. The good wear resistance of this alloy results from the high proportion of primary carbides that are formed on the basis of the high carbon content. However, a high proportion of primary carbides causes processing problems during the manufacture of this alloy as a wrought alloy.
For all mentioned alloys, the high-temperature strength or creep strength in the range of 500° C. to 900° C. is due to the additions of aluminum, titanium and/or niobium (or further elements such as Ta, etc.), which lead to precipitation of the γ′ and/or γ″ phase. Furthermore, the high-temperature strength or the creep strength is also improved by high contents of solid-solution-hardening elements such as chromium, aluminum, silicon, molybdenum and tungsten, as well as by a high carbon content.
Concerning the high-temperature corrosion resistance, it must be pointed out that alloys with a chromium content of around 20% form a chromium oxide layer (Cr2O3) that protects the material. In the course of service in the area of application, the chromium content is slowly consumed for buildup of the protective layer. Therefore the useful life of the material is improved by a higher chromium content, since a higher content of the element chromium forming the protective layer delays the point in time at which the Cr content falls below the critical limit and oxides other than Cr2O3 are formed, such as cobalt-containing and nickel-containing oxides, for example.
For processing of the alloy, especially during hot forming, it is necessary that no phases that greatly strain-harden the material, such as the γ′ or γ″ phase, for example, are formed at temperatures at which hot forming takes place, and thus lead to cracking during hot forming. At the same time, these temperatures must be sufficiently far below the solidus temperature of the alloy to prevent incipient melting in the alloy.
The task underlying the invention consists in conceiving a nickel-chromium wrought alloy that has
This task is accomplished by an age-hardening nickel-chromium-titanium-aluminum wrought alloy with very good wear resistance and at the same time very good high-temperature corrosion resistance, good creep strength and good processability, with (in mass-%) 25 to 35% chromium, 1.0 to 3.0% titanium, 0.6 to 2.0% aluminum, 0.005 to 0.10% carbon, 0.0005 to 0.050% nitrogen, 0.0005 to 0.030% phosphorus, max. 0.010% sulfur, max. 0.020% oxygen, max. 0.70% silicon, max. 2.0% manganese, max. 0.05% magnesium, max. 0.05% calcium, max. 2.0% molybdenum, max. 2.0% tungsten, max. 0.5% niobium, max. 0.5% copper, max. 0.5% vanadium, if necessary 0 to 20% Fe, if necessary 0 to 15% cobalt, if necessary 0 to 0.20% Zr, if necessary 0.0001 to 0.008% boron, the rest nickel and the usual process-related impurities, wherein the nickel content is greater than 35% and the following relationships must be satisfied:
Cr+Fe+Co≧26% (1)
in order to achieve good processability and
fh≧0 with (2a)
fh=6.49+3.88Ti+1.36Al−0.301Fe+(0.759−0.0209Co)Co−0.428Cr−28.2C (2)
in order that an adequate strength is achieved at higher temperatures, wherein Ti, Al, Fe, Co, Cr and C are the concentrations of the elements in question in mass-% and fh is expressed in %.
Advantageous improvements of the subject matter of the invention can be inferred from the associated dependent claims.
The variation range for the element chromium lies between 25 and 35%, wherein preferred ranges may be adjusted as follows:
The titanium content lies between 1.0 and 3.0%. Preferably Ti may be adjusted within the variation range as follows in the alloy:
The aluminum content lies between 0.6 and 2.0%, wherein here also, depending on service range of the alloy, preferred aluminum contents may be adjusted as follows:
The alloy contains 0.005 to 0.10% carbon. Preferably this may be adjusted within the variation range as follows in the alloy:
This is similarly true for the element nitrogen, which is contained in contents between 0.0005 and 0.05%. Preferred contents may be specified as follows:
The alloy further contains phosphorus in contents between 0.0005 and 0.030%. Preferred contents may be specified as follows:
The element sulfur is specified as follows in the alloy:
The element oxygen is contained in the alloy in contents of max. 0.020%. Preferred further contents may be specified as follows:
The element Si is contained in the alloy in contents of max. 0.70%. Preferred further contents may be specified as follows:
Furthermore, the element Mn is contained in the alloy in contents of max. 2.0%. Preferred further contents may be specified as follows:
The element Mg is contained in the alloy in contents of max. 0.05%. Preferred further contents may be specified as follows:
The element Ca is contained in the alloy in contents of max. 0.05%. Preferred further contents may be specified as follows:
The element niobium is contained in the alloy in contents of max. 0.5%. Preferred further contents may be specified as follows:
Molybdenum and tungsten are contained individually or in combination in the alloy with a content of maximum 2.0% each. Preferred contents may be specified as follows:
Furthermore, maximum 0.5% Cu may be contained in the alloy.
Beyond this, the content of copper may be limited as follows:
Furthermore, maximum 0.5% vanadium may be contained in the alloy.
Furthermore, the alloy may if necessary contain between 0.0 and 20.0% iron, which beyond this may be limited even more as follows:
Furthermore, the alloy may if necessary contain between 0.0 and 15% cobalt, wherein, depending on the area of application, preferred contents may be adjusted within the following variation ranges:
Furthermore, the alloy may if necessary contain between 0 and 0.20% zirconium, which beyond this may be limited even more as follows:
Furthermore, between 0.0001 and 0.008% boron may if necessary be contained in the alloy as follows. Preferred further contents may be specified as follows:
The nickel content should be higher than 35%. We may specify preferred further contents as follows:
The following relationship between Cr and Fe and Co must be satisfied in order to ensure an adequate resistance to wear:
Cr+Fe+Co≧26% (1)
wherein Cr, Fe and Co are the concentrations of the elements in question in mass-%.
Preferred further ranges may be adjusted with
Cr+Fe+Co≧27% (1a)
Cr+Fe+Co≧28% (1b)
Cr+Fe+Co≧29% (1c)
The following relationship between Ti, Al, Fe, Co, Cr and C must be satisfied in order that an adequately high strength at higher temperatures is achieved:
fh≧0 with (2a)
fh=6.49+3.88Ti+1.36Al−0.301Fe+(0.759−0.0209Co)Co−0.428Cr−28.2C (2)
wherein Ti, Al, Fe, Co, Cr and C are the concentrations of the elements in question in mass-% and fh is expressed in %.
Preferred ranges may be adjusted with
fh≧1% (2b)
fh≧3% (2c)
fh≧4% (2d)
fh≧5% (2e)
fh≧6% (2f)
fh≧7% (2f)
Optionally the following relationship between Cr, Mo, W, Fe, Co, Ti, Al and Nb may be satisfied in the alloy, in order that adequately good processability is achieved:
fver=≦7 with (3a)
fver=32.77+0.5932Cr+0.3642Mo+0.513W+(0.3123−0.0076Fe)Fe+(0.3351−0.003745Co−0.0109Fe)Co+40.67Ti*Al+33.28Al2−13.6TiAl2−22.99Ti−92.7Al+2.94Nb (3)
wherein Cr, Mo, W, Fe, Co, Ti, Al and Nb are the concentrations of the elements in question in mass-% and fver is expressed in %. Preferred ranges may be adjusted with
fver=≦5% (3b)
fver=≦3% (3c)
fver=<0% (3d)
Optionally the element yttrium may be adjusted in contents of 0.0 to 0.20% in the alloy. Preferably Y may be adjusted within the variation range as follows in the alloy:
Optionally the element lanthanum may be adjusted in contents of 0.0 to 0.20% in the alloy. Preferably La may be adjusted within the variation range as follows in the alloy:
Optionally the element Ce may be adjusted in contents of 0.0 to 0.20% in the alloy. Preferably Ce may be adjusted within the variation range as follows in the alloy:
Optionally, in the case of simultaneous addition of Ce and La, cerium mixed metal may also be used in contents of 0.0 to 0.20%. Preferably cerium mixed metal may be adjusted within the variation range as follows in the alloy:
Optionally 0.0 to 0.20% hafnium may also be contained in the alloy. Preferred ranges may be specified as follows:
Optionally 0.0 to 0.60% tantalum may also be contained in the alloy
Finally, the elements lead, zinc and tin may also be present as impurities in the following contents:
Pb max. 0.002%
Zn max. 0.002%
Sn max. 0.002%.
The alloy according to the invention is preferably melted in the vacuum induction furnace (VIM), but may also be melted under open conditions, followed by a treatment in a VOD or VLF system. After casting in ingots or possibly as continuous casting, the alloy is annealed if necessary at temperatures between 600° C. and 1100° C. for 0.1 to 100 hours, if necessary under protective gas such as argon or hydrogen, for example, followed by cooling in air or in the moving annealing atmosphere. Thereafter remelting may be carried out by means of VAR or ESR, if necessary followed by a 2nd remelting process by means of VAR or ESR. Then the ingots are annealed if necessary at temperatures between 900° C. and 1270° C. for 0.1 to 70 hours, then hot-formed, if necessary with one or more intermediate annealings between 900° C. and 1270° C. for 0.05 to 70 hours. The hot forming may be carried out, for example, by means of forging or hot rolling. Throughout the entire process, the surface of the material may if necessary be machined (even several times) intermediately and/or at the end chemically (e.g. by pickling) and/or mechanically (e.g. by cutting, by abrasive blasting or by grinding) in order to clean it. The control of the hot-forming process may be applied such that thereafter the semifinished product is already recrystallized with grain sizes between 5 and 100 μm, preferably between 5 and 40 μm. If necessary, solution annealing is then carried out in the temperature range of 700° C. to 1270° C. for 0.1 min to 70 hours, if necessary under protective gas such as argon or hydrogen, for example, followed by cooling in air, in the moving annealing atmosphere or in the water bath. After the end of hot forming, cold forming to the desired semifinished product form may if necessary (for example by rolling, drawing, hammering, stamping, pressing) with reduction ratios up to 98%, if necessary with intermediate annealings between 700° C. and 1270° C. for 0.1 min to 70 hours, if necessary under protective gas such as argon or hydrogen, for example, followed by cooling in air, in the moving annealing atmosphere or in the water bath. If necessary, chemical and/or mechanical (e.g. abrasive blasting, grinding, turning, scraping, brushing) cleanings of the material surface can be carried out intermediately in the cold-forming process and/or after the last annealing.
The alloys according to the invention or the finished parts made therefrom attain the final properties by age-hardening annealing between 600° C. and 900° C. for 0.1 to 300 hours, followed by cooling in air and/or in a furnace. By such an age-hardening annealing, the alloy according to the invention is age-hardened by precipitation of a finely dispersed γ′ phase. Alternatively, a two-stage annealing may also be carried out, wherein the first annealing takes place in the range of 800° C. to 900° C. for 0.1 to 300 hours, followed by cooling in air and/or furnace, and a second annealing takes place between 600° C. and 800° C. for 0.1 hours to 300 hours, followed by cooling in air.
The alloy according to the invention can be readily manufactured and used in the product forms of strip, sheet, rod, wire, longitudinally welded pipe and seamless pipe.
These product forms are manufactured with a mean grain size of 3 μm to 600 μm. The preferred range lies between 5 pμm and 70 μm, especially between 5 and 40 μm.
The alloy according to the invention can be readily processed by means of forging, upsetting, hot extrusion, hot rolling and similar processes. By means of these methods it is possible to manufacture components such as valves, hollow valves or bolts, among others.
It is intended that the alloy according to the invention will be used preferably in areas for valves, especially outlet valves of internal combustion engines. However, use in components of gas turbines, as fastening bolts, in springs and in turbochargers is also possible.
The parts manufactured from the alloy according to the invention, especially the valves or the valve seat faces, for example, may be subjected to further surface treatments, such a nitriding, for example, in order to increase the wear resistance further.
For measurement of the wear resistance, oscillating dry sliding wear tests were carried out in a pin-on-disk test bench (Optimol SRV IV tribometer). The radius of the hemispherical pins, which were polished to a mirror finish, was 5 mm. The pins were made from the material to be tested. The disk consisted of cast iron with a tempered, martensitic matrix with secondary carbides within a eutectic carbide network with the composition (C≈1.5%, Cr≈6%, S≈0.1%, Mn≈1%, Mo≈9%, Si≈1.5%, V≈3%, the rest Fe). The tests were carried out at a load of 20 N with a sliding path of one mm, a frequency of 20 Hz and a relative humidity of approximately 45% at various temperatures. Details of the tribometer and of the test procedure are described in C. Rynio, H. Hattendorf, J. Klöwer, H.-G. Lüdecke, G. Eggeler, Mat.-wiss. u. Werkstofftech. 44 (2013), 825. During the tests, the coefficient of friction, the linear displacement of the pin in disk direction (as a measure of the linear total wear of pin and disk) and the electrical contact resistance between pin and disk are continuously measured. Two different load-sensing modules, which are denoted in the following by (a) and (n), were used for the measurements. They yield results that are quantitatively slightly different but qualitatively similar. The load-sensing module (n) is the more accurate. After the end of a test, the volume loss of the pin was determined and used as a measure of the ranking for the wear resistance of the material of the pin.
The high-temperature strength was determined in a hot tension test according to DIN EN ISO 6892-2. For this purpose the offset yield strength Rp0.2 and the tensile strength Rm were determined. The tests were performed on round specimens with a diameter of 6 mm in the measurement area and an initial gauge length L0 of 30 mm. The specimens were taken transverse to the forming direction of the semifinished product. The crosshead speed for Rp0.2 was 8.33·10−5 l/s (0.5%/min) and for Rm was 8.33·10−4 l/s (5%/min).
The specimen was mounted at room temperature in a tension testing machine and heated to the desired temperature without being loaded with a tensile force. After the test temperature was reached, the specimen was maintained without load for one hour (600° C.) or two hours (700° C. to 1100° C.) for temperature equilibration. Thereafter the specimen was loaded with a tensile force such that the desired elongation rates were maintained and the test was begun.
The creep strength of a material is improved with increasing high-temperature strength. Therefore the high-temperature strength is also used for appraisal of the creep strength of the various materials.
The corrosion resistance at higher temperatures was determined in an oxidation test at 800° C. in air, wherein the test was interrupted every 96 hours and the changes in mass of the specimens due to the oxidation were determined. The specimens were confined in ceramic crucibles during the test, so that any oxide spalling off was collected, allowing the mass of spalled oxide to be determined by weighing the crucible containing the oxide. The sum of the mass of the spalled oxide and of the change in mass of the specimen is the gross change in mass of the specimen. The specific change in mass is the change in mass relative to the surface area of the specimens. In the following, these are denoted by mnet for the specific net change in mass, mgross for the specific gross change in mass and mspall for the specific change in mass of the spalled oxides. The tests were carried out on specimens with a thickness of approximately 5 mm. Three specimens were removed from each batch; the reported values are the mean values of these 3 specimens.
The phases occurring at equilibrium were calculated for the various alloy variants with the JMatPro program of Thermotech. The TTNI7 database for nickel-base alloys of Thermotech was used as the database for the calculations. In this way it is possible to identify phases that if formed embrittle the material in the service range. Furthermore, it is possible to identify the temperature ranges in which, for example, hot forming should not be carried out, since under those conditions phases form that greatly strain-harden the material and thus lead to cracking during hot forming. For a good processability, especially for hot forming, such as hot rolling, forging, upsetting, hot extrusion and similar processes, for example, an adequately broad temperature range in which such phases are not formed must be available.
In accordance with the stated task, the alloy according to the invention should have the following properties:
The new alloy should have a better wear resistance than the NiCr20TiAl reference alloy. Besides this material, Stellite 6 was also tested for comparison. Stellite 6 is a highly wear-resistant cobalt-base cast alloy with a network of tungsten carbides, consisting of approximately 28% Cr, 1% Si, 2% Fe, 6% W, 1.2% C, the rest Co, but because of its high carbide content it must be cast directly into the desired shape. By virtue of its network of tungsten carbides, Stellite 6 attains a very high hardness of 438 HV30, which is very advantageous for the wear. The alloy “E” according to the invention is supposed to approach the volume loss of Stellite 6 as closely as possible. The objective is in particular to decrease the high-temperature wear between 600 and 800° C., which is the relevant temperature range for application as an outlet valve, for example. Therefore the following criteria in particular should apply for the alloys “E” according to the invention:
Mean value of the volume loss (alloy “E”)≦0.50×mean value of the volume loss (NiCr20TiAl reference) at 600° C. or 800° C. (4a)
In the “low-temperature range” of the wear, the volume loss is not permitted to increase disproportionately. Therefore the following criteria should be additionally applicable.
Mean value of the volume loss (alloy “E”)≦1.30×mean value of the volume loss (NiCr20TiAl reference) at 25° C. and 300° C. (4b)
If a volume loss of NiCr20TiAl both for an industrial-scale batch and a reference laboratory batch is available in a series of measurements, the mean value of these two batches must be used in the inequalities (4a) and (4b).
Table 3 shows the lower end of the scatter band of the 0.2% offset yield strength for NiCr20TiAl in the age-hardened state at temperatures between 500 and 800° C., while Table 2 shows the lower end of the scatter band of the tensile strength.
The 0.2% offset yield strength of the alloy according to the invention should lie at least in this value range for 600° C. and should not be more than 50 MPa smaller than this value range for 800° C., in order to obtain adequate strength. This means in particular that the following values should be attained:
600° C.: Offset yield strength Rp0.2≧650 MPa (5a)
800° C.: Offset yield strength Rp0.2≧390 MPa (5b)
The inequalities (5a) and (5b) are attained in particular when the following relationship between Ti, Al, Fe, Co, Cr and C is satisfied:
fh≧0 with (2a)
fh=6.49+3.88Ti+1.36Al−0.301Fe+(0.759−0.0209Co)Co−0.428Cr−28.2C (2)
wherein Ti, Al, Fe, Co, Cr and C are the concentrations of the elements in question in mass-% and fh is expressed in %.
The alloy according to the invention should have a corrosion resistance in air similar to that of NiCr20TiAl.
For nickel-chromium-iron-titanium-aluminum alloys, the high-temperature strength or creep strength in the range of 500° C. to 900° C. depends on the additions of aluminum, titanium and/or niobium, which lead to precipitation of the γ′ and/or γ″ phase. If the hot forming of these alloys is carried out in the precipitation range of these phases, the risk of cracking exists. Thus the hot forming should preferably take place above the solvus temperature Tsγ′ (or Tsγ″) of these phases. To ensure that an adequate temperature range is available for the hot forming, the solvus temperature Tsγ′ (or Tsγ″) should be below 1020° C.
This is satisfied in particular when the following relationship between Cr, Mo, W, Fe, Co, Ti, Al and Nb is satisfied:
fver=≦7 with (3a)
fver=32.77+0.5932Cr+0.3642Mo+0.513W+(0.3123−0.0076Fe)Fe+(0.3351−0.003745Co−0.0109Fe)Co+40.67Ti*Al+33.28Al2−13.6TiAl2−22.99Ti−92.7Al+2.94Nb (3)
wherein Cr, Mo, W, Fe, Co, Ti, Al and Nb are the concentrations of the elements in question in mass-% and fver is expressed in %.
Tables 5a and 5b show the analyses of the batches melted on the laboratory scale together with some industrial-scale batches melted according to the prior art (NiCr20TiAl) and cited for reference. The batches according to the prior art are marked with a T, and those according to the invention with an E. The batches melted on the laboratory scale are marked with an L and the batches melted on the industrial scale with a G. Batch 250212 is NiCr20TiAl, but was melted at a laboratory batch and is used as reference.
The ingots of the alloys in Tables 5a and b melted on the laboratory scale in vacuum were annealed between 1100° C. and 125° C. for 0.1 to 70 hours and hot-rolled to a final thickness of 13 mm and 6 mm respectively by means of hot rolling and further intermediate annealings between 1100° C. and 1250° C. for 0.1 to 1 hour. The temperature control during hot rolling was such that the sheets were recrystallized. The specimens needed for the measurements were prepared from these sheets.
The comparison batches melted on an industrial scale were melted by means of VIM and cast as ingots. These ingots were remelted by ESR. These ingots were annealed between 1100° C. and 1250° C. for 0.1 min to 70 h, if necessary under protective gas such as argon or hydrogen, for example, followed by cooling in air, in the moving annealing atmosphere or in the water bath, and hot-rolled to a final diameter between 17 and 40 mm by means of hot rolling and further intermediate annealings between 1100° C. and 1250° C. for 0.1 to 20 hours. The temperature control during hot rolling was such that the sheets were recrystallized.
All alloy variants typically had a grain size of 21 to 52 μm (see Table 6).
After preparation of the specimens, these were age-hardened by an annealing at 850° C. for 4 hours/cooling in air followed by an annealing at 700° C. for 16 hours/cooling in air:
Table 6 shows the Vickers hardness HV30 before and after the age-hardening annealing. The hardness HV30 in the age-hardened state is in the range of 366 to 416 for all alloys except for batch 250330. Batch 250330 had a somewhat lower hardness of 346 HV30.
For the exemplary batches in Table 5a and 5b, the following properties are compared:
Wear tests were carried out at 25° C., 300° C., 600° C. and 800° C. on alloys according to the prior art and on the various laboratory heats. Most tests were repeated several times. Mean values and standard deviations were then determined.
The mean values±standard deviations of the measurements carried out are presented in Table 7. In the case of a single value, the standard deviation is missing. For orientation, the composition of the batches is roughly described in the alloy column of Table 7. In addition, the maximum values for the volume loss of the alloys according to the invention, from the inequalities (4a) for 600 and 800° C. respectively and (4b) for 25° C. and 300° C., are entered in the last row.
The volume losses at 600 and 800° C. are very small, and so the differences between various alloys can no longer be measured with certainty. Therefore a test was also carried out at 800° C. with 20 N for 2 hours+100 N for 5 hours, sliding path 1 mm, 20 Hz with load-sensing module (n), in order to cause a somewhat larger wear in the high-temperature range also. The results are plotted in
The comparison of the various alloys was performed at various temperatures. In
Especially on the basis of the values measured at 800° C., it was found that the volume loss of the pin in the wear test could be greatly reduced by a Cr content between 25 and 35% in the alloys according to the invention. Thus the batch 250326 according to the invention containing 30% Cr exhibits a reduction of the volume loss to 0.042±0.011% mm3 at 800° C. and to 0.026 mm3 even at 600° C., both smaller than or equal to 50% of the volume loss of NiCr20TiAl, the respective maximum value from (4a). At 300° C. the volume loss of 0.2588 mm3 was likewise below the maximum value from (4b), just as at 25° C., with 1.41±0.18 mm3 (load-sensing module (n)). Therefore chromium contents between 25 and 35% are of advantage in particular for wear at higher temperatures.
In the case of laboratory batch 250209 containing 10% Co, the volume loss at 800° C. decreased to 0.144±0.012 mm3, which is below the maximum value from (4a). At 25, 300 and 600° C., no increase of the wear was observed. In the case of laboratory batch 250329 containing 30% Co, the volume loss at 800° C. once again decreased significantly to 0.061±0.005 mm3, which is below the maximum value from (4a). The same was found at 600° C. with a decrease to 0.020 mm3, which is below the maximum value from (4a). At 25° C., the laboratory batch 250329 containing 30% Co exhibited a decrease to 0.93±0.02 mm3 with load-sensing module (n). Even at 300° C., this laboratory batch, with 0.244 mm3, exhibited a wear similar to that of reference batch 320776 and 250212, quite in contrast to the cobalt-base alloy Stellite 6, which at this temperature exhibited a significantly higher volume loss than reference batch 320776 and 250212. Thus the Co-containing laboratory batches satisfy the inequality (4a). Thus the optional addition of Co is advantageous. From cost viewpoints, a restriction of the optional content of cobalt to values between 0 and 15% is advantageous.
For laboratory batch 250330, a further reduction of the wear to 0.021±0.001 mm3 could be achieved by addition of 10% iron in addition to 29% Co. Thus an optional content of iron between 0 and 20% is advantageous.
For the volume losses measured at 800° C., it was found it was found on the basis of the laboratory batches 250325 (6.5% Fe), 250206 (11% Fe) and 250327 (29% Fe) that the volume loss of the pin in the wear test can be greatly reduced by an Fe content, such that is was smaller than or equal to 50% of the volume loss of NiCr20TiAl (4a) at one of the two temperatures, wherein the first % are particularly effective. Even at 25° C. and 300° C., the inequalities (4b) are satisfied by the alloys with an Fe content. Especially at 300° C., the alloys even had a volume loss reduced by more than 30%. Thus an optional content of iron between 0 and 20% is advantageous. An iron content also lowers the metal costs for this alloy.
In
The NiCr20TiAl alloys according to the prior art, batches 320776 and 250212, had a sum of Cr+Fe+Co equal to 20.3% and 20.2% respectively, both of which are smaller than 26%, and so did not meet the criteria (4a) and (4b) for a very good wear resistance, but especially not the criteria (4a) for a good high-temperature wear resistance. The batches 250211, 250214, 250208 and 250210 also did not meet the criteria for a good high-temperature resistance, especially (4a), and had a sum of Cr+Fe+Co equal to 20.4%, 20.2%, 20.3% and 20.3% respectively, all of which are smaller than 26%. The batches 250325, 250206, 250327, 250209, 250329, 250330 and 250326 with Fe and Co additions or with an increased Cr content, especially the batch 250326, met the criteria (4a) for 800° C., in some cases even additionally for 600° C., and had a sum of Cr+Fe+Co equal to 26.4%, 30.5%, 48.6%, 29.6%, 50.0%, 59.3% and 30.3% respectively, all of which are greater than 26%. Thus they satisfied Equation (1) for a very good wear resistance.
The offset yield strength Rp0.2 and the tensile strength Rm at room temperature (RT), 600° C. and 800° C. are presented in Table 8. The measured grain sizes and the values for fh are also presented. In addition, the minimum values from the inequalities (5a) and (5b) are entered in the last row.
At 600° C., as Table 8 shows, the offset yield strengths Rp0.2 of all laboratory batches (L), i.e. also of the batches (E) according to the invention, and of all industrial-scale batches (G) were greater than 650 MPa, and so criterion (5a) was met.
At 800° C., as Table 8 shows, the offset yield strengths Rp0.2 of all laboratory batches (L), i.e. also of the batches according to the invention, and of all industrial-scale batches (G) were greater than 390 MPa, and so inequality (5b) was satisfied.
A certain iron content in the alloy may be advantageous for cost reasons. Batch 250327 containing 29% Fe just satisfied this inequality (5b), since, as shown by the consideration of the laboratory batch 250212 (reference, similar to the industrial-scale batches, with Fe smaller than 3%) and also of the industrial-scale batches and of the batches 250325 (6.5% Fe), 250206 (11% Fe) and 250327 (29% Fe) according to the prior art, an increasing alloying content of Fe decreased the offset yield strength Rp0.2 in the tension test (see also
The consideration of the laboratory batch 250212 (reference, similar to the industrial-scale batches, without additions of Co) and also of the industrial-scale batches and of the batches 250209 (9.8% Co) and 250329 (30% Co) showed that a content of 9.8% Co increased the offset yield strength Rp0.2 in the tension test at 800° C. to 526 MPa, while a further increase to 30% Co led again to a slight decrease to 489 MPa (see also
The laboratory batch 250326 according to the invention showed that, with an addition of 30% Cr, the offset yield strength Rp0.2 in the tension test at 800° C. was reduced to 415 MPa, which was still well above the minimum value of 390 MPa. Therefore an alloying content of 35% Cr is regarded as the upper limit for the alloy according to the invention.
In
Table 9 shows the specific changes in mass after an oxidation test at 800° C. in air after 6 cycles of 96 h, i.e. a total of 576 h. The specific gross change in mass, the specific net change in mass and the specific change in mass of the spalled oxides after 576 h are presented in Table 9. The exemplary batches of the NiCr20TiAl alloys according to the prior art, batches 321426 and 250212, exhibited a specific gross change in mass of 9.69 and 10.84 g/m2 respectively and a specific net change in mass of 7.81 and 10.54 g/m2 respectively. Batch 321426 exhibited slight spalling. Batch 250326 with an increased Cr content of 30% according to the invention had a specific gross change in mass of 6.74 g/m2 and a specific net change in mass of 6.84 g/m2, which were below the range of the NiCr20TiAl reference alloys. The increase of the Cr content improves the corrosion resistance. Thus a Cr content of 25 to 35% is advantageous for the oxidation resistance of the alloy according to the invention.
The batches 250325 (Fe 6.5%), 250206 (Fe 11%) and 250327 (Fe 29%) exhibited a specific gross change in mass of 9.26 to 10.92 g/m2 and a specific net change in mass of 9.05 to 10.61 g/m2, which lie in the range of the NiCr20TiAl reference alloys. Thus an Fe content of up to 30% does not negatively influence the oxidation resistance. The Co-containing batches 250209 (Co 9.8%) and 250329 (Co 30%) also had a specific gross change in mass of 10.05 and 9.91 g/m2 respectively and a specific net change in mass of 9.81 and 9.71 g/m2 respectively, which likewise were in the range of the NiCr20TiAl reference alloys. The batch 250330 (29% Co, 10% Fe) behaved in just the same way, with a specific gross change in mass of 9.32 g/m- and a specific net change in mass of 8.98 g/m. Thus a Co content of up to 30% does not also negatively influence the oxidation resistance.
All alloys according to Table 5b contain Zr, which contributes as a reactive element to improvement of the corrosion resistance. Optionally, further reactive elements such as Y, La, Ce, cerium mixed metal, Hf, which improve the effectiveness in similar manner, may now be added.
The phase diagrams for the alloys in Table 5a and 5b were therefore calculated and the solvus temperature Tsγ′ was entered in Table 5a. The value for fver in accordance with Formula (3) was also calculated for the compositions in Tables 5a and 5b. fver is larger the higher the solvus temperature Tsγ′ is. All alloys in Table 5a, including the alloys according to the invention, have a calculated solvus temperature Tsγ′ lower than or equal to 1020° C. and meet criterion (3a): fver≦7%. The inequality fver≦7% (3a) is therefore a good criterion for obtaining an adequately broad hot-forming range and thus a good processability of the alloy.
The claimed limits for the alloys “E” according to the invention can be justified individually as follows:
Too low Cr contents mean that the Cr concentration sinks very quickly below the critical limit during use of the alloy in a corrosive atmosphere, and so a closed chromium oxide layer can no longer be formed. For an alloy with improved corrosion resistance, 25% is therefore the lower limit for chromium. Too high Cr contents raise the solvus temperature Tsγ′ too much, and so the processability is significantly impaired. Therefore 35% must be regarded as the upper limit.
Titanium increases the high-temperature resistance at temperatures in the range up to 900° C. by promoting the formation of the γ′ phase. In order to obtain an adequate strength, at least 1.0% is necessary. Too high titanium contents raise the solvus temperature Tsγ′ too much, and so the processability is significantly impaired. Therefore 3.0% must be regarded as the upper limit.
Aluminum increases the high-temperature resistance at temperatures in the range up to 900° C. by promoting the formation of the γ′ phase. In order to obtain an adequate strength, at least 0.6% is necessary. Too high aluminum contents raise the solvus temperature Tsγ′ too much, and so the processability is significantly impaired. Therefore 2.0% must be regarded as the upper limit.
Carbon improves the creep strength. A minimum content of 0.005% C is necessary for a good creep strength. Carbon is limited to maximum 0.10%, since at higher contents this element reduces the processability due to the excess formation of primary carbides.
A minimum content of 0.0005% N is necessary for cost reasons. N is limited to maximum 0.050%, since this element reduces the processability due to the formation of coarse carbonitrides.
The content of phosphorus should be lower than or equal to 0.030%, since this surface-active element impairs the oxidation resistance. A too-low phosphorus content increases the cost. The phosphorus content is therefore 0.0005%.
The contents of sulfur should be adjusted as low as possible, since this surface-active element impairs the oxidation resistance and the processability. Therefore max. 0.010% S is specified.
The oxygen content must be lower than or equal to 0.020%, in order to ensure manufacturability of the alloy.
Too high contents of silicon impair the processability. The Si content is therefore limited to 0.70%.
Manganese is limited to 2.0%, since this element reduces the oxidation resistance.
Even very low Mg contents and/or Ca contents improve the processing by the binding of sulfur, whereby the occurrence of low-melting NiS eutectics is prevented. At too high contents, intermetallic Ni—Mg phases or Ni—Ca phases may occur, which again significantly impair the processability. The Mg content or the Ca content is therefore limited respectively to maximum 0.05%.
Molybdenum is limited to max. 2.0%, since this element reduces the oxidation resistance.
Tungsten is limited to max. 2.0%, since this element likewise reduces the oxidation resistance and at the carbon contents possible in wrought alloys has no measurable positive effect on the wear resistance.
Niobium increases the high-temperature resistance. Higher contents increase the costs very greatly. The upper limit is therefore set at 0.5%.
Copper is limited to max. 0.5%, since this element reduces the oxidation resistance.
Vanadium is limited to max. 0.5%, since this element reduces the oxidation resistance.
Iron increases the wear resistance, especially in the high-temperature range. It also lowers the costs. It may therefore be present optionally between 0 and 20% in the alloy. Too high iron contents reduce the yield strength too much, especially at 800° C. Therefore 20% must be regarded as the upper limit.
Cobalt increases the wear resistance and the high-temperature strength/creep strength, especially in the high-temperature range. It also lowers the costs. It may therefore be present optionally between 0 and 20% in the alloy. Too high cobalt contents increase the costs too much. Therefore 20% must be regarded as the upper limit.
If necessary, the alloy may also contain Zr, in order to improve the high-temperature resistance and the oxidation resistance. For cost reasons, the upper limit is set at 0.20% Zr, since Zr is a rare element.
If necessary, boron may be added to the alloy, since boron improves the creep strength. Therefore a content of at least 0.0001% should be present. At the same time, this surface-active element impairs the oxidation resistance. Therefore max. 0.008% boron is specified.
Nickel stabilizes the austenitic matrix and is needed for formation of the γ′ phase, which contributes to the high-temperature strength/creep strength. At a nickel content below 35%, the high-temperature strength/creep strength is reduced too much, and so 35% is the lower limit.
The following relationship between Cr, Fe and Co must be satisfied, to ensure, as was explained in the examples, that an adequate wear resistance is achieved:
Cr+Fe+Co≧26% (1)
wherein Cr, Fe and Co are the concentrations of the elements in question in mass-%.
Furthermore, the following relationship must be satisfied, to ensure than an adequate strength at higher temperatures is achieved:
fh≧0 with (2a)
fh=6.49+3.88Ti+1.36Al−0.301Fe+(0.759−0.0209Co)Co−0.428Cr−28.2C (2)
wherein Ti, Al, Fe, Co, Cr and C are the concentrations of the elements in question in mass-% and fh is expressed in %. The limits for fh were justified in detail in the foregoing text.
If necessary, the oxidation resistance may be further improved with additions of oxygen-affine elements such as yttrium, lanthanum, cerium, hafnium. They do this by becoming incorporated in the oxide layer and blocking the diffusion paths of the oxygen at the grain boundaries therein.
For cost reasons, the upper limit of yttrium is defined as 0.20%, since yttrium is a rare element.
For cost reasons, the upper limit of lanthanum is defined as 0.20%, since lanthanum is a rare element.
For cost reasons, the upper limit of cerium is defined as 0.20%, since cerium is a rare element.
Instead of Ce and/or La, it is also possible to use cerium mixed metal. For cost reasons, the upper limit of cerium mixed metal is defined as 0.20%.
For cost reasons, the upper limit of hafnium is defined as 0.20%, since hafnium is a rare element.
If necessary, the ally may also contain tantalum, since tantalum also increases the high-temperature resistance by promoting the γ′ phase formation. Higher contents raise the costs very greatly, since tantalum is a rare element. The upper limit is therefore set at 0.60%.
Pb is limited to max. 0.002%, since this element reduces the oxidation resistance and the high-temperature resistance. The same applies for Zn and Sn.
Furthermore, the following relationship between Cr, Mo, W, Fe, Co, Ti, Al and Nb must be satisfied, to ensure that an adequate processability is achieved:
fver≦7 with (3a)
fver=32.77+0.5932Cr+0.3642Mo+0.513W+(0.3123−0.0076Fe)Fe+(0.3351−0.003745Co−0.0109Fe)Co+40.67TiAl+33.28Al2−13.6TiAl2−22.99Ti−92.7Al+2.94Nb (3)
wherein Cr, Mo, W, Fe, Co, Ti, Al and Nb are the concentrations of the elements in question in mass-% and fver is expressed in %. The limits for fh were justified in detail in the foregoing text.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10 2014 001 329.4 | Feb 2014 | DE | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/DE2015/000009 | 1/12/2015 | WO | 00 |
