Background Art Nickel-chromium-iron alloys are used primarily for their oxidation resistance and strength at elevated temperatures. Such alloys may be used, for example, as sheathing for electric heating elements and thermocouples.

Electrical resistance alloys are found in instruments and control equipment to measure and regulate electrical characteristics and in furnaces and appliances to generate heat. In the latter applications, elevated temperature characteristics are of prime importance. In common commercial terminology, electrical resistance alloys used for generation of heat are referred to as resistance heating alloys.

Resistance heating alloys such as CHROMEL-A (80 Ni; 20 Cr) are used in many varied applications--from small household appliances to large industrial furnaces. In appliances, resistance heating elements are designed for intermittent, short-term service at about 100 to 1090°C (about 212 to 2000°F).

In industrial furnaces, elements must often operate continuously at temperatures

as high as 1300°C (2350°F); as high as 1700°C (3100°F) for kilns used for firing ceramics; and occasionally as high as 2000°C (3600°F) for special applications.

The primary requirements of materials used for heating elements are high melting point, high electrical resistivity, reproducible temperature coefficient of resistance, good oxidation resistance in furnace environments, absence of volatile components, and resistance to contamination. Other desirable properties are elevated temperature creep strength, high emissivity, low thermal expansion and lower modulus (both of which help minimize thermal fatigue), resistance to thermal shock, and strength and ductility or workability at fabrication temperatures.

It is known that nickel-chromium-iron compositions are ductile alloys and thus are workable. They play an important role in heaters for the higher temperature ranges. Such heaters are constructed to provide more effective mechanical support for the heating element.

Oxidation resistance of nickel-chromium alloys at elevated temperatures is weakened by the limited adherence of an initial protective oxide layer to the base metal. The ability of an alloy to survive long exposures of 2000°F and above in air depends in large part on this protective layer remaining intact. Alloys presently used for electrical heating elements are made with carefully controlled techniques of melting, refinement and fabrication to maximize performance. However, variations in chemistry that occur between melts lead to differences in performance, as measured, for example, by accelerated life tests.

Various theories have been suggested to explain why some melts may show substantial improvement over the average performance. However, there has not been positive confirmation of any theory that has led to predictable improvement in life test results and consistently reproducible performance.

Illustrative of certain nickel-based chrome-iron-aluminum alloys are those disclosed in JP 59-85836. Such alloys, however, have a VHN above 500, which is antithetical to workability.

U. S. Patent No. 2,515,185 relates to nickel alloys, and more particularly to age hardenable nickel alloys. Such alloys, however, do not exhibit the requisite ductility. Nor does this patent envision the use of elements which promotes oxidation resistance of the nickel alloys. In addition, the precipitation of particles of secondary phases that promote hardening will adversely affect the resistivity of the base alloy, making it less desirable as an electrical heating element.

U. S. Patent No. 4,460,542, which is incorporated herein by reference, calls for the addition of yttrium to a chromium-aluminum-iron alloy which exhibits resistance to oxidation at high temperatures. However, this alloy only requires 14-18% chromium. U. S. Patent No. 4,671,931 abandons the concept of yttrium additions.

Disclosure Of The Invention It is accordingly an object of the present invention to provide a modified nickel-chromium-aluminum-iron alloy having superior resistance to oxidation at elevated temperatures.

It is additionally an object of the present invention to provide such an alloy which is characterized also by its workability.

The invention relates to an oxidation-resistant alloy for use in a high temperature or a thermal cycling environment. The alloy is relatively readily workable as it has a VHN below about 350. The alloy has a nickel-based matrix including 19-23% chromium, 3-6% aluminum, and 2-8% iron. All percentages herein are expressed as percent by weight unless otherwise specified. If a nickel- aluminum intermetallic compound, such as Ni3Al were precipitated (an ordered FCC compound referred to as"gamma prime"), the alloy would become strengthened (i. e. workability is lessened), a property which is antithetical to ductility.

To remove or stabilize oxygen and sulfur dissolved in the alloy, additions of carefully controlled quantities of calcium and yttrium are made to the melt so that about 0.005-0.05%, and preferably about 0.005-0.025% of calcium and about 0.01-0. 06%, and preferably about 0.02-0.06% of yttrium are present in the cast metal ingots. These residual dissolved elements enable the aluminum in the fabricated alloy to combine with oxygen in the environment to form a means

for healing a thermodynamically stable oxide scale if the layer is damaged or spalls during thermal cycling.

Brief Description Of The Drawings FIGURE 1 depicts the results of an accelerated life test of several alloys falling within the scope of the disclosed invention, wherein resistance change is plotted as a function of cycling time; and FIGURE 2 depicts the steps for processing the alloy of the present invention.

Best Mode (s) For Carrying Out The Invention The alloy of the present invention is an oxidation-resistant alloy for use in a high temperature, or thermal cycling environment that is relatively soft (having a Vickers hardness (hereinafter VHN) below about 500, and preferably below 350 for malleability) workable at room temperature. The alloy has a nickel- based matrix, including, by weight, about 19-23 chromium and about 3-6 aluminum. Also present is about 2-8, and preferably about 2-3 iron for inhibiting nucleation and growth of nickel aluminum intermetallic compounds which would otherwise adversely harden the alloy and cause local disturbance of a uniform distribution of aluminum.

The alloy also includes about 0.005-0.05, and preferably about 0.005-0.025 calcium and about and preferably about 0.02-0.06 yttrium for stabilizing oxygen or sulfur dissolved in the alloy, while retaining some calcium and yttrium in solid solution. As a result, the aluminum in the alloy may

combine with oxygen in the environment to form a self-healing means for repairing a thermodynamically stable oxide layer if it becomes damaged or spalls in use.

There is thus formed a self-healing, thermodynamically stable oxide layer disposed upon a surface of the alloy which is exposed to an oxidizing atmosphere over a range of temperatures for extended periods of time. The stable oxide layer protects the alloy from the oxidizing atmosphere.

Aluminum is added for oxidation resistance. Its favorable resistance to oxidation renders the alloy suitable for use in ceramic kilns and heat treating furnaces.

1. Comparison of Processing Steps With JP 59-85836 Following is a side-by-side comparison of the respective processing steps used in JP 59-85836 and those adopted by the inventors of the present application: TABLE 1

Step Japanese Method Hoskins Method (See Figure [JP 59-85836] [Invention] 2) A Make alloy samples Make alloy samples B Solution heat treatment Solution heat treatment 2 hr; air; 1200°C (2192°F) [about 3-12 hr; air; 2200°F C Hot roll to bar D Cool Air Cool E Age heat treatment [2 hr; air; Solution Heat treatment 650°C (1202°F) [about 2-12 hr; air; 2300°F F Hot roll to strip G Air cool Air cool H Annealed 1 hour; air; 2100°F I Water quenched; remove oxide to clean metal surface J Cold roll to strip K Annealed 1 hour; air; 2100°F L Water quenched; remove oxide to clean metal surface M Cold roll to strip N Anneal O Water quenched; remove oxide to clean metal surface P Cold roll to finish thickness; slit to width Q Bright annealing 1800°F or higher It is believed that the differences in the processing steps of the present invention with respect to the processing steps disclosed in JP 59-85836 are

at least partly responsible for the alloys of the present invention having a VHN (below about 350) which is substantially lower than the VHN attained by the alloys of JP 59-85836 (above about 500).

2. Comparative Testing The following experiments show that the wrought material produced by the inventors according to the chemistry specified in the tables below has a hardness below the VHN of 500 after following the solutioning and aging treatment outlined in JP 59-85836.

3. Effect of Solution and Aging Treatment On Hardness In the As-Cast Condition To evaluate the effect of solution and aging treatment on the hardness of the disclosed alloys in the as-cast condition, several melts were prepared having the compositions set forth in Table 2 below: TABLE 2<BR> Chemistry Compositions (Wt.%) for Alloy in Study Melt No. C S Al Si Ca Cr Fe Ni Y Zr 21427 .005 .0009 3.77 .07 .022 20.02 7.53 68.54 .04 <.01 21428 .006 .0010 3.73 .08 .001 20.25 5.18 70.71 .04 <.01 22158 .006 .0011 3.69 .08 .011 19.89 2.43 73.80 .05 .04 22159 .005 .0005 3.84 .04 .003 19.22 2.53 74.30 .02 .04 22161 .005 .0010 3.67 .08 .006 19.93 2.46 73.78 .02 .05 22282 .006 .0011 3.74 .06 .008 20.20 2.45 73.45 .04 .05 22283 .013 .0010 3.76 .08 .011 20.33 2.44 73.28 .04 .05

Table 3 illustrates the VHN of the as-cast material before solution treatment: TABLE 3 VHN of As-Cast Material Prior to Solution Treatment (1 kg Load Applied for 20 Seconds) Melt No. Trial #2Trial#3AverageTrial (VH) (VH) (VH) (VH) 23127026121427-1281 24127825921427-2258 25329327621428-1282 31331130921428-2304 29725128022158-1291 29626726922158-2243 378*31430322159-1291 28828828922159-2291 29030429922161-1303 27629228422161-2283 26926726822282-1267 26426726522282-2264 27627027122283-1266 31129630122283-2296 *-indicates a test result which was deemed to be erroneous and was not included in the calculations The samples identified with a"1"following the heat numbers were air-cooled following the solution treatment (2 hours at 1200°C). Table 4 details the hardness results obtained for each of these samples.

TABLE 4 VHN of Material Air-Cooled Following a 2 Hour Solution Treatment at 1200°C (1 kg Load Applied for 20 Seconds)

Melt No. Trial #1 Trial #2 Average Difference from (VH) (VH) (VH) As-Cast (i\VH) 21427-1 253 240 247-14 21428-1 241 251 246-30 22158-1 236 212 224-56 22159-1 257 240 249-54 22161-1 233 223 228-71 22282-1 223 227 225-43 22283-1 240 244 242-29 The samples identified with a"2"following the heat numbers were water quenched following the solutioning treatment (2 hours at 1200°C).

Table 5 details the hardness results obtained for each of these samples. The water-quench step produces material that is markedly softer than the air-cooled material. This indicates the probable formation of a second phase in the alloy at slower cooling rates.

TABLE 5 VHN of Material Water Quenched Following a 2 Hour Solution Treatment at 1200°C (1 kg Load Applied for 20 Seconds)

Melt No. Trial #1 Trial #2 Average Difference from (VH) (VH) (VH) As-Cast (IvVH) 190187-7321427-2183 179178-13221428-2176 22158-2 185 177 181-88 22159-2 179 184 182-108 22161-2 168 176 172-112 22282-2 167 130 149-117 22283-2 185 176 181-121 Table 6 details the hardness results obtained for each of the samples following the aging treatment (5 hours at 650°C and then air-cooled). Recall that after the initial solution treatment the samples followed by a"1"were air-cooled and the samples followed by a"2"were water quenched. The results show very little change in hardness from the as-cast condition. There also appears to be no correlation between the cooling rate after solution treatment and the hardness after aging.

TABLE 6<BR> VHN of As-Cast<BR> Material After Aging Treatment<BR> (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trial #3 Average Difference from As-Cast (#VH) (VH) (VH) (VH) (VH) 21427-1 232 251 251 245 -16 21427-2 262 260 251 258 -1 21428-1 299 268 233 267 -9 21428-2 250 258 233 247 -62 22158-1 266 241 269 259 -21 22158-2 250 246 260 252 -17 22159-1 300 284 313 299 -4 22159-2 243 238 265 249 -40 22161-1 275 294 297 289 -10 22161-2 269 248 247 255 -29 22282-1 284 277 272 278 10 22282-2 249 276 242 256 -9 22283-1 303 315 269 296 25 22283-2 216 258 240 238 -63

The results of the above studies show that the as-cast material the inventors produce as described in the appended claims has a VHN significantly lower than about 350-after following the solutioning and aging treatment outlined in JP 59-85836.

4. Effect of Solution and Aging Treatment on Hardness in the Wrought Condition In the following study, three samples (Melt Nos. 22230,22273 and 22283) were air induction melted, cast into ingots and hot rolled to strip. The strip was annealed one hour in air at 2100° followed by oxide removal to yield a clean metal surface. One sample (22270) was processed in the same manner as the above-mentioned three samples except that it has not yet been annealed in air.

Following is the chemistry composition for each of the heats used in this study.

TABLE 7<BR> Chemistry Compositions (Wt.%) for Alloy 126 Melts in Study Melt No. C S Al Si Ca Cr Fe Nl Y Zr 22230 .007 .0007 3.89 .03 .001 20.27 2.49 73.27 <.01 .04 22270 .016 .0011 3.58 .08 .001 19.96 2.47 73.81 .04 .04 22273 .011 .0010 3.72 .09 .007 20.32 2.48 73.27 .04 .06 22283 .013 .0010 3.76 .08 .011 20.33 2.44 73.28 .04 .05

Table 8 details the hardness results obtained for each of the samples in the hot worked condition prior to the solution treatment. Note the significantly higher hardness of the samples from heat 22270, presumably due to the fact that those samples had not been annealed.

TABLE 8 Initial VHN Of Hot Rolled Material Prior To Solution Treatment (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Average (VH) (VH) (VH) 29729722230-1297 29729722230-2297 22270-1386413400 38037422270-2388 29129022273-1289 22273-2 255 258 257 22283-1 278 288 283 22283-2 266 269 268 The samples identified with a"1"following the heat numbers were air-cooled following the solution treatment (2 hours at 1200°C).

Table 9 details the hardness results obtained for each of these samples. Note the fact that the hardness value for 22270-1 has fallen in line with the other three samples, indicating an effective reduction of residual stresses has taken place during the solution treatment for that sample.

TABLE 9 VHN of Material Air-Cooled Following a 2 Hour Solution Treatment at 1200°C (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Average Difference from (VH) (VH) (VH) Initial (AVH) 22230-1 258 276 267-30 258263-13722270-1268 269273-1722273-1276 22283-1 265 276 271-12 The samples identified with a"2"following the heat numbers were water quenched following the solution treatment (2 hours at 1200°C).

Table 10 details the hardness results obtained for each of these samples. The water-quench produces material that is markedly softer than the air- cooled material. This indicates the probable formation of a second phase in the alloy at slower cooling rates. Again note that the VHN value for 22270-2 is now comparable with the other samples.

TABLE 10 VHN of Material Water Quenched Following a 2 Hour Solution Treatment at 1200°C (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Average Difference from (VH) (VH) (VH) Initial (AVH) 22230-2 199 199 199-98 22270-2 194 193 194-180 22273-2 204 203 204-53 22283-2 227 218 223-45 Table 11 details the hardness results obtained for each of the samples following the aging treatment (5 hours at 650°C and then air-cooled).

Recall that after the initial solution treatment the samples followed by a"1"were air-cooled and the samples followed by a"2"were water quenched. The results show very little change in hardness from the initial hot-worked condition (except in the case of 22270, which had not been annealed). There also appears to be no correlation between the cooling rate after solution treatment and the hardness after aging.

TABLE 11<BR> VHN of Hot Rolled Material After Aging Treatment<BR> (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trial #3 Average Difference from Initial (#VH) (VH) (VH) (VH) (VH) 22230-1 300 288 296 295 -2 22230-2 267 262 258 262 -35 22270-1 274 288 265 276 -124 22270-2 236 239 245 240 -134 22273-1 278 286 269 278 -12 22273-2 280 271 285 272 +15 22283-1 285 340 317 314 +31 22283-2 254 258 263 258 -10

Again, the results of the study show that the wrought material prepared according to the disclosed regimen having the composition described in the appended claims has a hardness well below a hardness of about 350 following the solution treatment and aging treatment outlined in JP 59-85836.

5. Effect of Solution and Aging Treatment On The Hardness of Alloy Compositions Outside The Claimed Ranges In The As- Cast And Wrought Condition The significance of the claimed ranges of chemical composition are disclosed in the following study which illustrate the effect of solution and aging treatment on the hardness of alloy compositions falling outside the claimed range.

6. Effect of Solution and Aging Treatment On The Hardness of Compositions Falling Outside The Claimed Aluminum Ranges In The As-Cast and Wrought Conditions The earlier studies revealed that alloys falling within the claimed range do not produce material with a VHN of about 350 or more after the solutioning and aging treatment at the temperatures and times detailed in JP 59- 85836.

Selected compositions were produced by weighing up to 10 pound charges and melting them. Foundry and rolled samples were prepared. Foundry samples were analyzed for chemical composition and sectioned to prepare samples which represent the as-cast condition for study. The rolled samples were used to produce hot-work material to study the wrought condition of the compositions.

A. Procedure 1) Two charges of the inventive alloy were weighed up and melted in a 10-pound induction furnace. Foundry samples (approximately 1.25"1.25" x 8") and rolling samples (cast into Pyrex tubes approximately. 409"$ x 6") were obtained for each of the heats.

2) The foundry samples were sectioned and prepared for chemical analysis. In addition, two samples were obtained by sectioning a wafer into nearly equal half-circle disks approximately 0.25" thick. Two of these samples were obtained from each of the heats. The samples were identified using an engraver to avoid loss of identity.

3) The rolling samples were heated to 2000°F and rotary swaged from. 409"+ to. 326"+. The samples were then sectioned to obtain wafers approximately 0.25"thick. Two of these samples were obtained from each of the heats. The samples were identified using an engraver to avoid loss of identity.

4) A VHN test was performed on each of the samples.

5) The samples were placed in a furnace at 1200°C (2192°F) for 2 hours.

6) One cast sample and one wrought sample from each of the heats was water quenched. The other samples from each of the heats were air- cooled.

7) A VHN test was performed on each of the samples.

8) The samples were placed in a furnace at 650 ° C (1202 °F) for 5 hours.

9) The samples were removed from the furnace and allowed to air cool.

10) A VHN test was performed on each of the samples.

B. Discussion of Results Table 12 provides the analyzed chemistry composition for both of the heats used in the study.

TABLE 12<BR> Charge Aims And Chemistry Compositions (Wt.%) Melt No. C S Al Si Ca Cr Fe Ni Y Zr X1377 .024 .0015 6.20 .04 .001 22.57 2.31 68.82 .001 .03 X1378 .017 .0014 6.26 .03 .001 22.76 7.24 63.66 .001 .03

Table 13 details the hardness results obtained for each of the samples prior to the solution treatment.

TABLE 13 Initial VHN Of Material Prior To Solution Treatment (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trial #3 Average (VH) (VH) (VH) (AVH) X1377-1 392 336 364 (Cast) X1377-2 396 375 386 (Cast) X1377-1 339 350 338 342 (Wrought) X1377-2 308 307 315 310 (Wrought) X1378-1 353 334 344 (Cast) X 1378-2 316 337 327 (Cast) X1378-1 332 333 330 332 (Wrought) X1378-2 330 338 336 335 (Wrought) The samples identified with a"1"following the heat numbers were air-cooled following the solution treatment (2 hours at 1200°C). Table 14 details the hardness results obtained for each of these samples.

TABLE 14<BR> VHN Of Material Air-Cooled Following<BR> A 2 Hour Solution Treatment At 1200°C<BR> (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trial #3 Average Difference from Initial (# (VH) (VH) (VH) (VH) X1377-1 342 340 341 -23 (Cast) X1377-1 337 327 324 329 -13 (Wrought) X1378-1 313 331 322 -22 (Cast) X1378-1 318 303 311 311 -21 (Wrought)

The samples identified with a"2"following the heat numbers were water quenched following the solution treatment (2 hours at 1200°C).

Table 15 details the hardness results obtained for each of these samples. The water-quench produces material that is markedly softer than the air- cooled material.

TABLE 15<BR> VHN Of Material Water Quenched Following<BR> A 2 Hour Solution Treatment At 1200°C<BR> (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trial #3 Average Difference from Initial (#VH) (VH) (VH) (VH) (VH) X1377-2 337 324 331 -55 (Cast) X1377-2 267 263 267 266 -44 (Wrought) X1378-2 252 256 254 -73 (Cast) X1378-2 251 249 245 248 -87 (Wrought

Table 16 details the hardness results obtained for each of the samples following the aging treatment (5 hours at 650°C and then air-cooled).

Recall that after the initial solution treatment the samples followed by a"1"were air-cooled and the samples followed by a"2"were water quenched. There is a correlation between the aging treatment and hardenability in these compositions, especially in the high Al, Cr and Fe sample (X1378). In all cases, the effect was strongest on samples that were water-quenched after the solutioning treatment.

TABLE 16<BR> VHN Of Material After Aging Treatment<BR> (1 kg Load Applied for 20 Seconds) Melt No. Trial #1 Trial #2 Trials #3 Trial #4 Average Difference from Initial (VH) (VH) (VH) (VH) (VH) X1377-1 363 369 349 360 -4 (Cast) X1377-2 495 424 386 435 +49 (Cast) X1377-1 363 347 354 348 353 +11 (Wrought) X1377-2 360 360 366 362 +52 (Wrought) X1378-1 340 351 353 348 +4 (Cast) X1378-2 546 544 544 545 +218 (Cast) X1378-1 384 399 453 401 409 +77 (Wrought) X1378-2 519 489 569 502 520 +185 (Wrought)

The study shows that, for example, the composition produced in Heat X1378 (High Al, Cr, and Fe) has a VHN greater than 500 following the solution and aging treatment outlined in JP 59-85836. Only the material that was water quenched after solution treatment produced material with a VHN of at least 500.

7. Microstructural Comparison Of Alloys Of The Present Invention With Respect To Alloys Of JP 59-85836 To compare the microstructure of the alloys of the present invention to the alloys disclosed in JP 59-85836, the following five alloys were prepared: TABLE 17<BR> Compositions Of The Examined Alloys<BR> Elements (Wt.%) Alloy No. C S Al Si Ca Cr Fe Ni Y Zr 22283 0.013 0.0010 3.76 0.08 0.011 20.33 2.44 73.28 0.04 0.05 X1377-2a 6.26 22.57 2.30 b X1378-2a 6.26 22.76 7.24 b 17032c 3.95 36.5 8.0 b 17033d 3.96 38.5 5.8 b aOther elements are assumed to be similar to alloy 22283.<BR> bBalance<BR> cExample 3 in JP 59-85836<BR> dComp. Example 2 in JP 59-85836

The 22283, X1377-2 and X1378 alloys were prepared in accordance with the present invention. Alloy 22283 is a more preferred composition of the present invention. Alloy X1377-2 has amounts of aluminum and chromium near the upper limit of the present invention. Alloy X1378-2 has amounts of aluminum, chromium and iron near the upper limits of the present invention.

All of the alloys were evaluated in the cast condition. The final heat treatment was a solution treatment for 2 hr. at 1200°C followed by water- quenching. The quenched samples were aged for 5 hr at 650°C.

All of the samples were evaluated for their microhardness. The data are listed below in Table 18.

TABLE 18 Microhardness Data Alloy No. Microhardness (HV) 22283-1 272.8 + 6.2 X1377-2 337.8 + 71.9 X1378-2 518. 1 + 98.98. 17032 656.5 + 23.4 17033 653.2 + 18.4 Microstructures of the alloys were examined and reported in MICROSTRUCTURAL EVALUATION OF NICKEL-BASED SAMPLES FROM HOSKINS MANUFACTURING COMPANY, by V. K. Sikka dated March 1998, which is incorporated herein by reference. The following observations are detailed in the above report.

Alloy X1378-2 This alloy shows two distinctly different phases. High- magnification scanning electron microscopy (SEM) shows the two phases to be distinctly different. The second phase has a fairly large fraction of the microstructure.

Alloy X1377-2 This alloy also shows distinct second phase, but the amount is much smaller than Alloy X1378-2.

Alloys 17032 and 17033 These alloys show second phase distributed throughout. High-magnification SEMs of alloy X1378-2 show that the second phase has a very fine lamellar microstructure. A careful microhardness measurement (see Table 19) shows that the matrix has a hardness of 288 + 21 and that the lamellar phase a value of 655 9. These large differences in hardness of the phases are responsible for a very large deviation in average hardness, when no distinction between the two is made during typical VHN measurement.

TABLE 19 Selective Microhardness Data Of Various Samples

Alloy No. Microhardnessa X1377-2 Matrix 331 + 4 Second phase 540 + 44 X1378-2 Matrix 288 21 Second phase 655 9 17032 700 14 17033 705 9 22283-1 289 + 14 'VHN at 500 g load.

SEMs micrographs of alloy X1377-2 show the presence of very fine, nearly spherical particles inside the second-phase regions. It is believed that the lamellar structure in alloy 1378-2 has changed its morphology to spherical particles in alloy X1377-2. The morphology change probably is a direct result of the difference in the iron content of the two alloys. The careful microhardness measurements of the matrix and the second-phase region in alloy X1377-2 again show the second phase to be much harder (540 44 VHN) as opposed to the matrix (331 + 44 VHN).

It is believed that the larger standard deviation in the hardness of the second phase in alloy X1377-2 may be due to nonuniformity in distribution of the spherical particles as opposed to perfectly aligned particles in the second phase of alloy X1378-2.

A microprobe analysis in Table 20 indicates the presence of large gamma-prime particles at the grain boundaries. It is believed that the lamellar looking microstructure is also gamma prime, but was too fine for chemical analysis. It is believed that spherical particles in alloy X1377-2 are the same as the lamellar structure in alloy 1378-2, and that they are gamma prime.

Micrographs of alloys 17032 and 17033 show fairly uniform distribution of large particles rather than distinct two phase regions observed in alloys X1377-2 and X1378-2. The repeated microhardness of alloys 17032 and 17033 gave values of 700 14 and 705 9, respectively. The uniformity of the hardness in these samples suggest that the hardening phase is uniformly distributed. Higher magnification micrographs show the presence of a very fine lamellar structure in both alloys 17032 and 17033. This lamellar structure is somehow different than the very uniformly-spaced lamellar structure in alloy X1378-2.

TABLE 20 Microprobe Analysis Summary

Alloy No. Phasea 22283-1 No second phase. X1377-2 No second phase. X1378-2 Grey phase" (56.92 Ni, 8.94 Cr, 3.54 Fe, 30.59 Al, at %) 17032 Grey phase' (27.16 Ni, 62.13 Cr, 7.40 Fe, 3.31 Al, at %) 17033 Grey phased (21.04 Ni, 71.07 Cr, 5.12 Fe, 2.54 Al, 0.23 Ti, at %)

aKey phases, when observed, were quantitatively analyzed. bThis composition is enriched in aluminum and is typical of gamma prime.

Typically, the gamma prime is Ni3 (Al, Ti, Nb, Ta). It could be (Ni Cr Fe) 3Al.

This phase is rich in chromium and very low in aluminum. Its composition matches that of alpha-chromium. dThis phase is very rich in chromium and very low in aluminum, and contains a small amount of titanium. Its composition matches that of alpha- chromium.

The microprobe analysis in Table 20 had shown the coarse gray particles to be alpha-chromium. It is believed that the same alpha-chromium particles that are present as finely spaced lamellar caused the large hardening observed in alloys 17032 and 17033.

The titanium addition in alloy 17033 did not seem to have a significant effect in modifying the microstructure.

Our observations of gamma-chromium hardening in alloys 17032 and 17033 are consistent with that proposed in the Japanese Patent No. 59-85836.

The nominal alloy 22283-1 did not show any features in the optical structure. Its second check of microhardness of 289+/-14 VHN is consistent essentially with a solid solution alloy.

A combination of microprobe, microhardness, optical microstructure, and high-magnification SEM analysis of five different specimens results in the following observations: 1. Alloys X1378-2 and X1377-2 consist of two phases (the matrix and the second phase). The second phase had significantly higher (655 and 540, respectively) hardness than the matrix phase (288 and 331, respectively).

The second phase hardening is believed to be from very fine gamma prime. The gamma prime in the low iron content, alloy X1377-2 is typical of gamma prime (spherical particles) observed in nickel-based superalloys. However, the higher iron containing alloy X1378-2 contained gamma prime as the lamellar morphology;

2. The duplicated sample from the Japanese Patent No. 59- 85836, alloys 17032 and 17033, showed large hardening from a very fine distribution of alpha-chromium particles. These results are consistent with the hardening mechanisms proposed from the Japanese Patent No. 59-85836; and 3. The standard alloy 22283 was observed to be a solid solution. No second phase was observed in this alloy.

8. Additional Examples To further illustrate various aspects of the present invention, the following additional examples are disclosed.

Seven alloys were melted. Their chemistry, aside from impurities, is summarized in Table 21: TABLE 21 Nominal Starting Chemical Composition (Wt. %)

Alloy No. FeAlYCaCr 1Bal.62- 3.78NA20.01320.22 62-2 Bal. 19.55 2. 27 3. 73 NA 0.0109 3Bal.19.334.283.67NA0.010962- 64-1 Bal. 20.50 0. 15 4. 10 NA 0.012 64-2 Bal. 20.23 2. 30 4. 00 0. 012 0.0063 66-1 Bal. 19. 90 6. 00 4. 20 None 0.0160 Detected 66-2 Bal. 19. 40 7. 80 4. 10 0. 040 0.0060 'Compositions are of the solid alloy.

2"Not Added." Oxidation tests were conducted at 2200°F for 80 hours and at 2300°F for 195 hours to compare oxidation resistance of the seven alloys.

The experimental procedure involved placing samples of the alloys in an electrically heated box furnace and exposing them to air flow created by convective thermal currents. The samples were cycled daily. After cooling to room temperature during each cycle, the specimens were examined. The results appear in Table 22: TABLE 22 Typical Elevated Temperature Properties

Oxidation Wt. Change Alloy No. (Gram/Cm2) ASTM Life (Hours) 1-62- 0.00156 120 2-62- 0.00153 212 3-62- 0.001918 96 64-1 0.000860 224 64-2 0.000437 800 66-1 0.000112 444 66-2 0.000226 582 In Table 22, a specific weight loss is caused by spalling of the oxide scale during cycling. However, Table 22 also shows that following an effective reduction in the sulfur content by combination with calcium and/or yttrium, the alloy may exhibit a slight weight gain through the formation of a protective scale of alumina. Desulfurized specimens tend to show a positive specific weight change throughout the duration of the life test.

The right-hand column of Table 22 follows the procedure set forth in ASTM designation B76-90 standard test method (2 min. on/2 min. off) at 2150°F for determining the accelerated life of nickel-chromium and nickel- chromium-iron alloys used for electrical heating. That publication is incorporated herein by reference.

Results obtained after 80 hours at 2200°F and 195 hours at 2300°F (31 cycles to room temperature) indicate that alloys such as 66-1 exhibited preferable performance because they incurred the lowest weight gain (0.000112

gm/cm2). However, alloys 64-2 and 66-2, which contain yttrium, yielded the longest life (800 hours and 582 hours, respectively).

Although lifetime improvement over conventional alloys is not susceptible of easy quantification, it is likely that lifetime is increased by at least several fold.

Turning now to Figure 1, there is depicted the results of an accelerated life test wherein resistance change (percent) is plotted against hours of cycling. The 64-2 alloy, as noted above, lasted almost 800 hours.

Table 23 summarizes the disclosed and preferred ranges of alloy composition, aside from impurities: TABLE 23 Alloy Composition Ranges3 (Wt. %) ELEMENT PREFERRED Ni Balance Balance Cr 19-23 19-21 Fe 2-8 2-3 Al 3-6 3.5-4.5 Ca 0.005-0.05 0.005-0.025 0.02-0.06Y0.01-0.06 3 Compositions are of the solid alloy.

Without wishing to be bound by any particular theory, the inventors have observed that the advanced alloys and coatings disclosed rely on the formation and adherence of a thin and continuous aluminum oxide film to protect the base alloy from further oxidation attack at elevated temperatures. In order for the alumina scale to serve its protective function, it must remain adherent to the

underlying alloy under prolonged exposure and thermal cycling conditions. It is known that segregation of indigenous sulfur to the metal-oxide interface induces premature scale spalling of the scale. This may occur through a reduction in the interfacial adhesion strength, with a resulting reduction in component lifetime.

The inherent reactivity of yttrium requires an exceptionally high degree of control over alloy chemistry during melting/casting. Control of the concentration of the reactive element additions is particularly important, since retention of a minimum amount in solution in cast alloys is required to impart acceptable oxidation resistance. However, when the concentration of the reactive element greatly exceeds that of the impurities with which it reacts, the formation of extraneous, low melting point phases can result. If the proportion of the reactive element is too high or too low, the oxidation characteristics of the alloy may be suboptimal.

The inventors have discovered that by the addition of aluminum to the nickel-chromium base alloy, oxidation resistance is enhanced by the formation of an impervious layer of aluminum oxide. To ensure that the surface oxide layer remains intact with a"self-healing"mechanism if the oxide is damaged or spalls, it is necessary to have aluminum dissolved uniformly in the alloy matrix to a level of about 3-6%, and preferably about 4% by weight. At this level, it is thought that diffusion of the aluminum atoms in the matrix can occur quickly to replace aluminum depletion by alumina formation at the surface. The limited solubility of aluminum in the nickel-chromium alloy can result in precipitation of some of the aluminum in the form of a nickel-aluminum phase referred to as"gamma prime."As noted earlier, these particles can cause severe hardening in the alloy and a reduction in the aluminum in solution in the surrounding matrix.

To modulate this mechanism, iron and other metals may be added.

To enhance the adherence of the surface oxide to the base metal, small additions of calcium and yttrium with other rare earth metals have been found to be

effective. The results of alloying and testing show that reproducible results can be assured if the following chemistry controls are exercised: TABLE 24 Detailed Composition Ranges (Wt. %) Chromium 19-23 Aluminum 3-6 Iron 2-8 Zirconium 0-0.10 Calcium 0.005-0.05 Yttrium 0. 01-0. 06 Sulfur 0-0.008 Boron 0-0. 005 -Carbon0 0.03 -Silicon0 2.0 Manganese 0-0.5 Titanium 0-0.25 Nickel Balance The inventors have developed the disclosed alloy so that it may serve as a thermocouple sheath, tube, wire, or strip for use as a heating element or as a tubular member in applications which are exposed to an oxidizing atmosphere at high temperatures.

The disclosed alloys provide for favorable oxidation resistance at the highest temperatures of intended use without spalling of the surface oxide.

Yttrium, calcium, and zirconium in the proper relative amounts effectively reduce the oxygen and sulfur content of the resulting alloy. As a result, these highly

reactive additions are present uniformly in the matrix of the alloy. This ensures initial reaction of aluminum in the matrix at the hot surface with the ambient air/atmosphere and provides a base for bonding subsequent layers of aluminum oxide. The control of sulfur in the alloy by calcium and yttrium tends to neutralize the potential of this undesirable impurity to interfere with oxide layer formation.

The chemistry of the alloys disclosed requires refining the melt to neutralize the sulfur and oxygen contained in the alloy. Hot fabricability is promoted through the addition of zirconium in the disclosed amounts and sequence during melting. The precipitation of"gamma prime"is retarded by increasing the solid solubility of aluminum in the nickel-chromium base alloy. As a result, cold working processes are facilitated.

Further, the disclosed alloys promote formation of a continuous protective layer of alumina. Thus, adherence of the oxide layer to the base alloy is ensured during thermal cycling, thereby promoting longer life at higher temperatures over comparable alloys which are presently available.

It will be apparent to those skilled in the art that the novel principles of the invention disclosed herein, in connection with specific examples thereof, will support various other modifications and applications of the same. It is accordingly desired that, in construing the breadth of the appended claims, they will not be limited to the specific examples of the invention described herein.