US 5,205,875 AGrant
Wrought Gamma Titanium Aluminide Alloys Modified by Chromium, Boron, and Nionium
Issue Date:1993-04-27
•5 Claims
•3 Drawing Sheets
Abstract
A TiAl composition is prepared to have high strength and to have improved ductility by altering the atomic ratio of the titanium and aluminum to have what has been found to be an effective aluminum concentration and by addition of chromium, boron, and niobium according to the approximate formula Ti-Al.sub.46-48 Cr.sub.2 Nb.sub.2 B.sub.0.1-0.2. The composition is preferably prepared by casting, homogenization at a high temperature, and forging the homogenized casting.
Metadata
Assignee
- General Electric Company
Inventor
- Shyh-Chin Huang
Application Information
Application Number:US 8015570
Filing Date:1991-12-02
Priority Date:1991-12-02
Art Unit:111
Classifications
IPC:
C22C 1400
Field of Search:
420148418;421421;670
Patent Drawings (3 sheets)
Description
CROSS-REFERENCE TO RELATED APPLICATIONS
The subject application relates to copending applications as follows:
Ser. No. 07/812393, filed Dec. 23, 1991, Ser. No. 07/801556, filed Dec. 2, 1991, Ser. No. 07/801558, filed Dec. 2, 1991, and Ser. No. 07/811371, filed Dec. 20, 1991.
Ser. No. 07/354,965, filed May 22, 1989; Ser. Nos. 07/546,962, and 07/546,973, both filed Jul. 2, 1990; Ser. Nos. 07/589,823, and 07/589,827, both filed Sep. 26, 1990; Ser. No. 07/613,494, filed Jun. 12, 1991; Ser. Nos. 07/631,988, and 07/631,989, both filed Dec. 21, 1990; Ser. No. 07/695,043, filed May 2, 1991; and Ser. No. 07/739,004, filed Aug. 1, 1991.
The text of these related applications are incorporated herein by reference.
BACKGROUND OF THE INVENTION
The present invention relates generally to alloys of titanium and aluminum. More particularly, it relates to gamma alloys of titanium and aluminum which have been modified both with respect to stoichiometric ratio and with respect to chromium, boron, and niobium addition.
It is known that as aluminum is added to titanium metal in greater and greater proportions the crystal form of the resultant titanium aluminum composition changes. Small percentages of aluminum go into solid solution in titanium and the crystal form remains that of alpha titanium. At higher concentrations of aluminum (including about 25 to 35 atomic %) an intermetallic compound Ti.sub.3 Al is formed. The Ti.sub.3 Al has an ordered hexagonal crystal form called alpha-2. At still higher concentrations of aluminum (including the range of 50 to 60 atomic % aluminum) another intermetallic compound, TiAl, is formed having an ordered tetragonal crystal form called gamma.
The alloy of titanium and aluminum having a gamma crystal form, and a stoichiometric ratio of approximately one, is an intermetallic compound having a high modulus, a low density, a high thermal conductivity, favorable oxidation resistance, and good creep resistance. The relationship between the modulus and temperature for TiAl compounds to other alloys of titanium and in relation to nickel base superalloys is shown in FIG. 3. As is evident from the figure, the TiAl has the best modulus of any of the titanium alloys. Not only is the TiAl modulus higher at higher temperature but the rate of decrease of the modulus with temperature increase is lower for TiAl than for the other titanium alloys. Moreover, the TiAl retains a useful modulus at temperatures above those at which the other titanium alloys become useless. Alloys which are based on the TiAl intermetallic compound are attractive lightweight materials for use where high modulus is required at high temperatures and where good environmental protection is also required. The present invention relates to improvements in the gamma titanium aluminides.
One of the characteristics of TiAl which limits its actual application to such uses is a brittleness which is found to occur at room temperature. Also, the strength of the intermetallic compound at room temperature needs improvement before the TiAl intermetallic compound can be exploited in structural component applications. Improvements of the TiAl intermetallic compound to enhance ductility and/or strength at room temperature are very highly desirable in order to permit use of the compositions at the higher temperatures for which they are most suitable.
With potential benefits of use at light weight and at high temperatures, what is most desired in the TiAl compositions which are to be used is a combination of strength and ductility at room temperature. A minimum ductility of the order of one percent is acceptable for some applications of the metal composition but higher ductilities are much more desirable. A minimum strength for a composition to be useful is about 50 ksi or about 350 MPa. However, materials having this level of strength are of marginal utility and higher strengths are often preferred for some applications.
The stoichiometric ratio Of TiAl compounds can vary over a range without altering the crystal structure. The aluminum content can vary from about 50 to about 60 atom percent. The properties of TiAl compositions are subject to very significant changes as a result of relatively small changes of one percent or more in the stoichiometric ratio of the titanium and aluminum ingredients. Also, the properties are similarly affected by the addition of similar relatively small amounts of ternary elements.
I have now discovered that further improvements can be made in the gamma TiAl intermetallic compounds by incorporating therein a combination of additive elements so that the composition not only contains a ternary additive element but also a quaternary additive element and a dopant.
The additive elements are chromium and niobium, and the dopant is boron.
Furthermore, I have discovered that the composition including the quaternary additive element and dopant has a uniquely desirable combination of properties which include a desirably high ductility and a valuable oxidation resistance.
PRIOR ART
There is extensive literature on the compositions of titanium aluminum including the Ti.sub.3 Al intermetallic compound, the gamma TiAl intermetallic compounds and the Ti.sub.3 Al intermetallic compound. A patent, U.S. Pat. No. 4,294,615, entitled "Titanium Alloys of the TiAl Type" contains an extensive discussion of the titanium aluminide type alloys including the gamma TiAl intermetallic compound. As is pointed out in the patent in column 1, starting at line 50, in discussing TiAl's advantages and disadvantages relative to Ti.sub.3 Al:
"It should be evident that the TiAl gamma alloy system has the potential for being lighter inasmuch as it contains more aluminum. Laboratory work in the 1950's indicated that titanium aluminide alloys had the potential for high temperature use to about 1000.degree. C. But subsequent engineering experience with such alloys was that, while they had the requisite high temperature strength, they had little or no ductility at room and moderate temperatures, i.e., from 20.degree. to 550.degree. C. Materials which are too brittle cannot be readily fabricated, nor can they withstand infrequent but inevitable minor service damage without cracking and subsequent failure. They are not useful engineering materials to replace other base alloys."
It is known that the alloy system TiAl is substantially different from Ti.sub.3 Al (as well as from solid solution alloys of Ti) although both TiAl and Ti.sub.3 Al are basically ordered titanium aluminum intermetallic compounds. As the '615 patent points out at the bottom of column 1:
"Those well skilled recognize that there is a substantial difference between the two ordered phases. Alloying and transformational behavior of Ti.sub.3 Al resemble those of titanium, as the hexagonal crystal structures are very similar. However, the compound TiAl has a tetragonal arrangement of atoms and thus rather different alloying characteristics. Such a distinction is often not recognized in the earlier literature."
The '615 patent does describe the alloying of TiAl with vanadium and carbon to achieve some property improvements in the resulting alloy.
The '615 patent also discloses in Table 2 alloy T.sub.2 A-112 which is a composition in atomic percent of Ti-45Al-5.0 Nb but the patent does not describe the composition as having any beneficial properties.
A number of technical publications dealing with the titanium aluminum compounds as well as with characteristics of these compounds are as follows:
1. E. S. Bumps, H. D. Kessler, and M. Hansen, "Titanium-Aluminum System", Journal of Metals, TRANSACTIONS AIME, Vol. 194 (June 1952) pp. 609-614,
2. H. R. Ogden, D. J. Maykuth, W. L. Finlay, and R. I. Jaffee, "Mechanical Properties of High Purity Ti-Al Alloys", Journal of Metals, TRANSACTIONS AIME, Vol. 197 (February, 1953) pp. 267-272.
3. Joseph B. McAndrew and H. D. Kessler, "Ti-36 Pct Al as a Base for High Temperature Alloys", Journal of Metals, TRANSACTIONS AIME, Vol. 206 (October 1956) pp. 1345-1353.
4. S. M. Barinov, T. T. Nartova, Yu L. Krasulin and T. V. Mogutova, "Temperature Dependence of the Strength and Fracture Toughness of Titanium Aluminum", Izv. Akad. Nauk SSSR, Met., Vol. 5 (1983) p. 170.
In reference 4, Table I, a composition of titanium-36 aluminum -0.01 boron is reported and this composition is reported to have an improved ductility. This composition corresponds in atomic percent to Ti.sub.50 Al.sub.49.97 B.sub.0.03.
5. S. M. L. Sastry, and H. A. Lispitt, "Plastic Deformation of TiAl and Ti.sub.3 Al", Titanium 80 (Published by American Society for Metals, Warrendale, Pa.), Vol. 2 (1980) page 1231.
6. Patrick L. Martin, Madan G. Mendiratta, and Harry A. Lispitt, "Creep Deformation of TiAl and TiAl+W Alloys", Metallurgical Transactions A, Vol. 14A (October 1983) pp. 2171-2174.
7. Tokuzo Tsujimoto, "Research, Development, and Prospects of TiAl Intermetallic Compound Alloys", Titanium and Zirconium, Vol. 33, No. 3, 159 (July 1985) pp. 1-13.
8. H. A. Lispitt, "Titanium Aluminides--An Overview", Mat. Res. Soc. Symposium Proc., Materials Research Society, Vol. 39 (1985) pp. 351-364.
9. S. H. Whang et al., "Effect of Rapid Solidification in Ll.sub.o TiAl Compound Alloys", ASM Symposium Proceedings on Enhanced Properties in Struc. Metals Via Rapid Solidification, Materials Week (October 1986) pp. 1-7.
10. Izvestiya Akademii Nauk SSR, Metally. No. 3 (1984) pp 164-168.
11. P. L. Martin, H. A. Lispitt, N. T. Nuhfer and J. C. Williams, "The Effects of Alloying on the Microstructure and Properties of Ti.sub.3 Al and TiAl", Titanium 80 (published by the American Society of Metals, Warrendale, Pa.), Vol. 2 (1980) pp. 1245-1254.
12. D. E. Larsen, M. L. Adams, S. L. Kampe, L. Christodoulou, and J. D. Bryant, "Influence of Matrix Phase Morphology on Fracture Toughness in a Discontinuously Reinforced XD.TM. Titanium Aluminide Composite", Scripta Metallurgica et Materialia, Vol. 24, (1990) pp. 851-856.
13. Akademii Nauk Ukrain SSR, Metallofiyikay No. 50 (1974).
14. J. D. Bryant, L. Christodon, and J. R. Maisano, "Effect of TiB.sub.2 Additions on the Colony Size of Near Gamma Titanium Aluminides", Scripta Metallurgica et Materialia, Vol. 24 (1990) pp. 33-38.
The McAndrew reference discloses work under way toward development of a TiAl intermetallic gamma alloy. In Table II, McAndrew reports alloys having ultimate tensile strength of between 33 and 49 ksi as adequate "where designed stresses would be well below this level". This statement appears immediately above Table II. In the paragraph above Table IV, McAndrew states that tantalum, silver and (niobium) columbium have been found useful alloys in inducing the formation of thin protective oxides on alloys exposed to temperatures of up to 1200.degree. C. FIG. 4 of McAndrew is a plot of the depth of oxidation against the nominal weight percent of niobium exposed to still air at 1200.degree. C. for 96 hours. Just above the summary on page 1353, a sample of titanium alloy containing 7 weight % columbium (niobium) is reported to have displayed a 50% higher rupture stress properties than the TiAl used for comparison.
Commonly owned patents relating to gamma titanium aluminides include U.S. Pat. Nos. 4,842,817, 4,842,819, 4,836,983; 4,857,268; 4,879,092; 4,897,127; 4,902,474; 4,923,534; 5,028,491; 5,032,357; and 5,045,406.
A number of other patents also deal with TiAl compositions as follows:
U.S. Pat. No. 3,203,794 to Jaffee discloses various TiAl compositions.
Canadian Patent 621884 to Jaffee similarly discloses various compositions of TiAl.
U.S. Pat. No. 4,661,316 (Hashimoto) teaches titanium aluminide compositions which contain various additives.
Commonly owned U.S. Pat. No. 4,916,028 concerns a gamma TiAl alloy containing chromium, niobium, and carbon.
U.S. Pat. No. 4,842,820, assigned to the same assignee as the subject application, teaches the incorporation of boron to form a tertiary TiAl composition and to improve ductility and strength.
U.S. Pat. No. 4,639,281 to Sastry teaches inclusion of fibrous dispersoids of boron, carbon, nitrogen, and mixtures thereof or mixtures thereof with silicon in a titanium base alloy including Ti-Al.
European patent application 0275391 to Nishiyama teaches TiAl compositions containing up to 0.3 weight percent boron and 0.3 weight percent boron when nickel and silicon are present. No niobium is taught to be present in a combination with boron.
U.S. Pat. No. 4,774,052 to Nagle concerns a method of incorporating a ceramic, including boride, in a matrix by means of an exothermic reaction to impart a second phase material to a matrix material including titanium aluminides.
BRIEF STATEMENT OF THE INVENTION
In one of its broader aspects, the objects of the present invention are achieved by providing a nonstoichiometric TiAl base alloy, and adding a relatively low concentration of chromium and a low concentration of niobium as well as a boron dopant to the nonstoichiometric composition.
Addition of chromium in the order of approximately 1 to 3 atomic percent and of niobium to the extent of 1 to 5 atomic percent and boron to the extent of 0.1 to 0.3 atomic percent is contemplated.
The alloy of this invention may also be produced in wrought ingot form and may be processed by ingot metallurgy.
BRIEF DESCRIPTION OF THE DRAWINGS
The detailed description of the invention which follows will be understood with greater clarity if reference is made to the accompanying drawings in which:
FIG. 1 is a graph displaying ductility in relation to temperature of heat treatment.
FIG. 2 is a graph illustrating the relationship between load in pounds and crosshead displacement in mils for TiAl compositions of different stoichiometry tested in 4-point bending.
FIG. 3 is a graph illustrating the relationship between modulus and temperature for an assortment of alloys.
DETAILED DESCRIPTION OF THE INVENTION
There are a series of background and current studies which led to the findings on which the present invention involving the combined addition of chromium, niobium, and boron to a gamma TiAl are based. The first 25 examples deal with the background studies and the later examples deal with the current studies.
EXAMPLES 1-3
Three individual melts were prepared to contain titanium and aluminum in various stoichiometric ratios approximating that of TiAl. The compositions, annealing temperatures and test results of tests made on the compositions are set forth in Table I.
For each example, the alloy was first made into an ingot by electro arc melting. The ingot was processed into ribbon by melt spinning in a partial pressure of argon. In both stages of the melting, a water-cooled copper hearth was used as the container for the melt in order to avoid undesirable melt-container reactions. Also, care was used to avoid exposure of the hot metal to oxygen because of the strong affinity of titanium for oxygen.
The rapidly solidified ribbon was packed into a steel can which was evacuated and then sealed. The can was then hot isostatically pressed (HIPed) at 950.degree. C. (1740.degree. F.) for 3 hours under a pressure of 30 ksi. The HIPing can was machined off the consolidated ribbon plug. The HIPed sample was a plug about one inch in diameter and three inches long.
The plug was placed axially into a center opening of a billet and sealed therein. The billet was heated to 975.degree. C. (1787.degree. F.) and was extruded through a die to give a reduction ratio of about 7 to 1. The extruded plug was removed from the billet and was heat treated.
The extruded samples were then annealed at temperatures as indicated in Table I for two hours. The annealing was followed by aging at 1000.degree. C. for two hours. Specimens were machined to the dimension of 1.5.times.3.times.25.4 mm (0.060.times.0.120.times.1.0 in.) for four point bending tests at room temperature. The bending tests were carried out in a 4-point bending fixture having an inner span of 10 mm (0.4 in.) and an outer span of 20 mm (0.8 in.). The load-crosshead displacement curves were recorded. Based on the curves developed, the following properties are defined:
(1) Yield strength is the flow stress at a cross head displacement of one thousandth of an inch. This amount of cross head displacement is taken as the first evidence of plastic deformation and the transition from elastic deformation to plastic deformation. The measurement of yield and/or fracture strength by conventional compression or tension methods tends to give results which are lower than the results obtained by four point bending as carried out in making the measurements reported herein. The higher levels of the results from four point bending measurements should be kept in mind when comparing these values to values obtained by the conventional compression or tension methods. However, the comparison of measurements' results in many of the examples herein is between four point bending tests, and for all samples measured by this technique, such comparisons are quite valid in establishing the differences in strength properties resulting from differences in composition or in processing of the compositions.
(2) Fracture strength is the stress to fracture.
(3) Outer fiber strain is the quantity of 9.71 hd, where "h" is the specimen thickness in inches, and "d" is the cross head displacement of fracture in inches. Metallurgically, the value calculated represents the amount of plastic deformation experienced at the outer surface of the bending specimen at the time of fracture.
The results are listed in the following Table I. Table I contains data on the properties of samples annealed at 1300.degree. C. and further data on these samples in particular is given in FIG. 2.
It is evident from the data of this Table that alloy 12 for Example 2 exhibited the best combination of properties. This confirms that the properties of Ti-Al compositions are very sensitive to the Ti/Al atomic ratios and to the heat treatment applied. Alloy 12 was selected as the base alloy for further property improvements based on further experiments which were performed as described below.
It is also evident that the anneal at temperatures between 1250.degree. C. and 1350.degree. C. results in the test specimens having desirable levels of yield strength, fracture strength and outer fiber strain. However, the anneal at 1400.degree. C. results in a test specimen having a significantly lower yield strength (about 20% lower); lower fracture strength (about 30% lower) and lower ductility (about 78% lower) than a test specimen annealed at 1350.degree. C. The sharp decline in properties is due to a dramatic change in microstructure due, in turn, to an extensive beta transformation at temperatures appreciably above 1350.degree. C.
EXAMPLES 4-13
Ten additional individual melts were prepared to contain titanium and aluminum in designated atomic ratios as well as additives in relatively small atomic percents.
Each of the samples was prepared as described above with reference to Examples 1-3.
The compositions, annealing temperatures, and test results of tests made on the compositions are set forth in Table II in comparison to alloy 12 as the base alloy for this comparison.
For Examples 4 and 5, heat treated at 1200.degree. C., the yield strength was unmeasurable as the ductility was found to be essentially nil. For the specimen of Example 5 which was annealed at 1300.degree. C., the ductility increased, but it was still undesirably low.
For Example 6, the same was true for the test specimen annealed at 1250.degree. C. For the specimens of Example 6 which were annealed at 1300.degree. and 1350.degree. C. the ductility was significant but the yield strength was low.
None of the test specimens of the other Examples were found to have any significant level of ductility.
It is evident from the results listed in Table II that the sets of parameters involved in preparing compositions for testing are quite complex and interrelated. One parameter is the atomic ratio of the titanium relative to that of aluminum. From the data plotted in FIG. 3, it is evident that the stoichiometric ratio or nonstoichiometric ratio has a strong influence on the test properties which formed for different compositions.
Another set of parameters is the additive chosen to be included into the basic TiAl composition. A first parameter of this set concerns whether a particular additive acts as a substituent for titanium or for aluminum. A specific metal may act in either fashion and there is no simple rule by which it can be determined which role an additive will play. The significance of this parameter is evident if we consider addition of some atomic percentage of additive X.
If X acts as a titanium substituent, then a composition Ti.sub.48 Al.sub.48 X.sub.4 will give an effective aluminum concentration of 48 atomic percent and an effective titanium concentration of 52 atomic percent.
If, by contrast, the X additive acts as an aluminum substituent, then the resultant composition will have an effective aluminum concentration of 52 percent and an effective titanium concentration of 48 atomic percent.
Accordingly, the nature of the substitution which takes place is very important but is also highly unpredictable.
Another parameter of this set is the concentration of the additive.
Still another parameter evident from Table II is the annealing temperature. The annealing temperature which produces the best strength properties for one additive can be seen to be different for a different additive. This can be seen by comparing the results set forth in Example 6 with those set forth in Example 7.
In addition, there may be a combined concentration and annealing effect for the additive so that optimum property enhancement, if any enhancement is found, can occur at a certain combination of additive concentration and annealing temperature so that higher and lower concentrations and/or annealing temperatures are less effective in providing a desired property improvement.
The content of Table II makes clear that the results obtainable from addition of a ternary element to a nonstoichiometric TiAl composition are highly unpredictable and that most test results are unsuccessful with respect to ductility or strength or to both.
EXAMPLES 14-17
A further parameter of the gamma titanium aluminide alloys which include additives is that combinations of additives do not necessarily result in additive combinations of the individual advantages resulting from the individual and separate inclusion of the same additives.
Four additional TiAl based samples were prepared as described above with reference to Examples 1-3 to contain individual additions of vanadium, niobium, and tantalum as listed in Table III. Two of these compositions are the optimum compositions reported in commonly owned U.S. Pat. Nos. 4,842,817, and 4,857,268.
The fourth composition is a composition which combines the vanadium, niobium and tantalum into a single alloy designated in Table III to be alloy 48.
From Table III, it is evident that the individual additions vanadium, niobium and tantalum are able on an individual basis in Examples 14, 15, and 16 to each lend substantial improvement to the base TiAl alloy. However, these same additives when combined into a single combination alloy do not result in a combination of the individual improvements in an additive fashion. Quite the reverse is the case.
In the first place, the alloy 48 which was annealed at the 1350.degree. C. temperature used in annealing the individual alloys was found to result in production of such a brittle material that it fractured during machining to prepare test specimens.
Secondly, the results which are obtained for the combined additive alloy annealed at 1250.degree. C. are very inferior to those which are obtained for the separate alloys containing the individual additives.
In particular, with reference to the ductility, it is evident that the vanadium was very successful in substantially improving the ductility in the alloy 14 of Example 14. However, when the vanadium is combined with the other additives in alloy 48 of Example 17, the ductility improvement which might have been achieved is not achieved at all. In fact, the ductility of the base alloy is reduced to a value of 0.1.
Further, with reference to the oxidation resistance, the niobium additive of alloy 40 clearly shows a very substantial improvement in the 4 mg/cm2 weight loss of alloy 40 as compared to the 31 mg/cm2 weight loss of the base alloy. The test of oxidation, and the complementary test of oxidation resistance, involves heating a sample to be tested at a temperature of 982.degree. C. for a period of 48 hours. After the sample has cooled, it is scraped to remove any oxide scale. By weighing the sample both before and after the heating and scraping, a weight difference can be determined. Weight loss is determined in mg/cm2 by dividing the total weight loss in grams by the surface area of the specimen in square centimeters. This oxidation test is the one used for all measurements of oxidation or oxidation resistance as set forth in this application.
For the alloy 60 with the tantalum additive, the weight loss for a sample annealed at 1325.degree. C. was determined to be 2 mg/cm2 and this is again compared to the 31 mg/cm2 weight loss for the base alloy. In other words, on an individual additive basis both niobium and tantalum additives were very effective in improving oxidation resistance of the base alloy.
However, as is evident from Example 17, results listed in Table III alloy 48 which contained all three additives, vanadium, niobium and tantalum in combination, the oxidation is increased to about double that of the base alloy. This is seven times greater than alloy 40 which contained the niobium additive alone and about 15 times greater than alloy 60 which contained the tantalum additive alone.
The individual advantages or disadvantages which result from the use of individual additives repeat reliably as these additives are used individually over and over again. However, when additives are used in combination the effect of an additive in the combination in a base alloy can be quite different from the effect of the additive when sued individually and separately int eh same base alloy. Thus, it has been discovered that addition of vanadium is beneficial to the ductibility of titanium aluminum compositions and this is disclosed and discussed in the commonly owned U.S. Pat. No. 4,857,268. Further, one of the additives which has been found to be beneficial to the strength of the TiAl base is the additive niobium. It has been shown by the McAndrew paper discussed above that the individual addition of niobium additive to TiAl base alloy can improve oxidation resistance. Similarly, the individual addition of tantalum is taught by McAndrew as assisting in improving oxidation resistance. Furthermore, in commonly owned U.S. Pat. No. 4,842,817, it is disclosed that addition of tantalum results in improvements in ductility.
In other words, it has been found that vanadium can individually contribute advantageous ductility improvements to gamma titanium aluminum compound and that tantalum can individually contribute to ductility and oxidation improvements. It has been found separately that niobium additives can contribute beneficially to the strength and oxidation resistance properties of titanium aluminum. However, the Applicant has found, as is indicated from this Example 17, that when vanadium, tantalum, and niobium are used together and are combined as additives in an alloy composition, the alloy composition is not benefited by the additions but rather there is a net decrease or loss in properties of the TiAl which contains the niobium, the tantalum, and the vanadium additives. This is evident from Table III.
From this, it is evident that, while it may seem that if two or more additive elements individually improve TiAl that their use together should render further improvements to the TiAl, it is found, nevertheless, that such additions are highly unpredictable and that, in fact, for the combined additions of vanadium, niobium and tantalum a net loss of properties result from the combined use of the combined additives together rather than resulting in some combined beneficial overall gain of properties.
However, from Table III above, it is evident that the alloy containing the combination of the vanadium, niobium and tantalum additions has far worse Oxidation resistance than the base TiAl 12 alloy of Example 2. Here, again, the combined inclusion of additives which improve a property on a separate and individual basis have been found to result in a net loss in the very property which is improved when the additives are included on a separate and individual basis.
EXAMPLES 18 thru 23
Six additional samples were prepared as described above with reference to Examples 1-3 to contain chromium modified titanium aluminide having compositions respectively as listed in Table IV.
Table IV summarizes the bend test results on all of the alloys, both standard and modified, under the various heat treatment conditions deemed relevant.
The results listed in Table IV offer further evidence of the criticality of a combination of factors in determining the effects of alloying additions or doping additions on the properties imparted to a base alloy. For example, the alloy 80 shows a good set of properties for a 2 atomic percent addition of chromium. One might expect further improvement from further chromium addition. However, the addition of 4 atomic percent chromium to alloys having three different TiAl atomic ratios demonstrates that the increase in concentration of an additive found to be beneficial at lower concentrations does not follow the simple reasoning that if some is good, more must be better. And, in fact, for the chromium additive just the opposite is true and demonstrates that where some is good, more is bad.
As is evident from Table IV, each of the alloys 49, 79 and 88, which contain "more" (4 atomic percent) chromium shows inferior strength and also inferior outer fiber strain (ductility) compared with the base alloy.
By contrast, alloy 38 of Example 18 contains 2 atomic percent of additive and shows only slightly reduced strength but greatly improved ductility. Also, it can be observed that the measured outer fiber strain of alloy 38 varied significantly with the heat treatment conditions. A remarkable increase in the outer fiber strain was achieved by annealing at 1250.degree. C. Reduced strain was observed when annealing at higher temperatures. Similar improvements were observed for alloy 80 which also contained only 2 atomic percent of additive although the annealing temperature was 1300.degree. C. for the highest ductility achieved.
For Example 20, alloy 87 employed the level of 2 atomic percent of chromium but the concentration of aluminum is increased to 50 atomic percent. The higher aluminum concentration leads to a small reduction in the ductility from the ductility measured for the two percent chromium compositions with aluminum in the 46 to 48 atomic percent range. For alloy 87, the optimum heat treatment temperature was found to be about 1350.degree. C.
From Examples 18, 19 and 20, which each contained 2 atomic percent additive, it was observed that the optimum annealing temperature increased with increasing aluminum concentration.
From this data it was determined that alloy 38 which has been heat treated at 1250.degree. C., had the best combination of room temperature properties. Note that the optimum annealing temperature for alloy 38 with 46 at. % aluminum was 1250.degree. C. but the optimum for alloy 80 with 48 at. % aluminum was 1300.degree. C.
These remarkable increases in the ductility of alloy 38 on treatment at 1250.degree. C. and of alloy 80 on heat treatment at 1300.degree. C. were unexpected as is explained in the commonly owned U.S. Pat. No. 4,842,819.
What is clear from the data contained in Table IV is that the modification of TiAl compositions to improve the properties of the compositions is a very complex and unpredictable undertaking. For example, it is evident that chromium at 2 atomic percent level does very substantially increase the ductility of the composition where the atomic ratio of TiAl is in an appropriate range and where the temperature of annealing of the composition is in an appropriate range for the chromium additions. It is also clear from the data of Table IV that, although one might expect greater effect in improving properties by increasing the level of additive, just the reverse is the case because the increase in ductility which is achieved at the 2 atomic percent level is reversed and lost when the chromium is increased to the 4 atomic percent level. Further, it is clear that the 4 percent level is not effective in improving the TiAl properties even though a substantial variation is made in the atomic ratio of the titanium to the aluminum and a substantial range of annealing temperatures is employed in studying the testing the change in properties which attend the addition of the higher concentration of the additive.
EXAMPLE 24
Samples of alloys were prepared which had a composition as follows:
Test samples of the alloy were prepared by two different preparation modes or methods and the properties of each sample were measured by tensile testing. The methods used and results obtained are listed in Table V immediately below.
In Table V, the results are listed for alloy samples 38 which were prepared according to two Examples, 18' and 24, which employed two different and distinct alloy preparation methods in order to form the alloy of the respective examples. In addition, test methods were employed for the metal specimens prepared from the alloy 38 of Example 18' and separately for alloy 38 of Example 24 which are different from the test methods used for the specimens of the previous examples. Turning now first to Example 18', the alloy of this example was prepared by the method set forth above with reference to Examples 1-3. This is a rapid solidification and consolidation method. In addition for Example 18', the testing was not done according to the 4 point bending test which is used for all of the other data reported in the tables above and particularly for Example 18 of Table IV above. Rather the testing method employed was a more conventional tensile testing according to which a metal samples are prepared as tensile bars and subjected to a pulling tensile test until the metal elongates and eventually breaks. For example, again with reference to Example 18' of Table V, the alloy 38 was prepared into tensile bars and the tensile bars were subjected to a tensile force until there was a yield or extension of the bar at 93 ksi.
The yield strength in ksi of Example 18' of Table V, measured by a tensile bar, compares to the yield strength in ksi of Example 18 of Table IV which was measured by the 4 point bending test. In general, in metallurgical practice, the yield strength determined by tensile bar elongation is a more generally used and more generally accepted measure for engineering purposes.
Similarly, the tensile strength in ksi of 108 represents the strength at which the tensile bar of Example 18' of Table V broke as a result of the pulling. This measure is referenced to the fracture strength in ksi for Example 18 in Table IV. It is evident that the two different tests result in two different measures for all of the data.
With regard next to the plastic elongation, here again there is a correlation between the results which are determined by 4 point bending tests as set forth in Table IV above for Example 18 and the plastic elongation in percent set forth in the last column of Table V for Example 18'.
Referring again now to Table V, the Example 24 is indicated under the heading "Processing Method" to be prepared by cast and forge ingot metallurgy. As used herein, the term "cast and forge ingot metallurgy" refers to a first step melting of the ingredients of the alloy 38 in the proportions set forth in Table V and corresponding exactly to the proportions set forth for Example 18'. In other words, the composition of alloy 38 for both Example 18' and for Example 24 are identically the same. (They are also exactly the same for alloy 38 of Example 18 of Table IV.)
The difference between the two examples of Table V is that the alloy of Example 18' was prepared by rapid solidification and the alloy of Example 24 was prepared by cast and forge ingot metallurgy. Again, the cast and forge ingot metallurgy involves a melting of the ingredients and solidification of the ingredients into an ingot followed by a forging of the cast ingot. The rapid solidification method involves the formation of a ribbon by the melt spinning method followed by the consolidation of the ribbon into a fully dense coherent metal sample.
In the cast and forge ingot processing procedure of Example 24 the ingot was prepared to a dimension of about 2" in diameter and about 1/2" thick in the approximate shape of a hockey puck. Following the melting and solidification of the hockey puck-shaped ingot, the ingot was enclosed within a steel annulus having a wall thickness of about 1/2" and having a vertical thickness which matched identically that of the hockey puck-shaped ingot. Before being enclosed within the retaining ring the hockey puck ingot was homogenized by being heated to 1250.degree. C. for two hours. The assembly of the hockey puck and containing ring were heated to a temperature of about 975.degree. C. The heated sample and containing ring were forged to a thickness of approximately half that of the original thickness. This procedure is referred to herein as a cast and forge processing.
Following the forging and cooling of the specimen, tensile specimens were prepared corresponding to the tensile specimens prepared for Example 18'. These tensile specimens were subjected to the same conventional tensile testing as was employed in Example 18' and the yield strength, tensile strength and plastic elongation measurements resulting from these tests are listed in Table V for Example 24. As is evident from the Table V results, the individual test samples were subjected to different annealing temperatures prior to performing the actual tensile tests.
For Example 18' of Table V, the annealing temperature employed on the tensile test specimen was 1250.degree. C. For the three samples of the alloy 38 of Example 24 of Table V, the samples were individually annealed at the three different temperatures listed in Table V and specifically 1225.degree. C., 1250.degree. C., and 1275.degree. C. Following this annealing treatment for approximately two hours, the samples were subjected to conventional tensile testing and the results again are listed in Table V for the three separately treated tensile test specimens.
Turning now again to the test results which are listed in Table V, it is evident that the yield strengths determined for the rapidly solidified alloy are somewhat higher than those which are determined for the ingot processed metal specimens. Also, it is evident that the plastic elongation of the samples prepared through the cast and forge ingot metallurgy route have generally higher ductility than those which are prepared by the rapid solidification route. The results listed for Example 24 demonstrate that although the yield strength measurements are somewhat lower than those of Example 18' they are fully adequate for many applications in aircraft engines and in other industrial uses. However, based on the ductility measurements and the results of the measurements as listed in Table 24 the gain in ductility makes the alloy 38 as prepared through the ingot metallurgy route a very desirable and unique alloy for those applications which require a higher ductility. Generally speaking, it is well-known that processing by ingot metallurgy is far less expensive than processing through melt spinning or rapid solidifications inasmuch as there is no need for the expensive melt spinning step itself nor for the consolidation step which must follow the metal spinning.
EXAMPLE 25
Samples of an alloy containing both chromium additive and niobium additive were prepared as disclosed above with reference to Examples 1-3. Tests were conducted on the samples and the results are listed in Table VI immediately below.
The data in Table VI evidences that unique properties are found in the gamma titanium aluminide containing both chromium and niobium. This unique composition is the subject of commonly owned U.S. Pat. No. 4,879,092.
EXAMPLES 26-29
Four additional samples of alloys were prepared according to the ingot metallurgy procedure set forth in Example 24 above. This set of four alloys was prepared by a cast and HIP procedure. The cast and HIP procedure involves first preparing a melt of the alloy to be cast and then casting the alloy into an ingot. The ingot is cut into bars or pins which can be conveniently subjected to a HIPing operation by enclosing each pin in a metal wrap and subjecting the wrap and its contents to a pressure of about 45 ksi at a temperature of about 1,050.degree. C.
Sample alloys were prepared according to this cast and HIP procedure and the conventional tensile properties of the alloys as prepared were tested. The test results are presented in Table VII immediately below.
Referring now to the contents of Table VII, the Example 2B is a binary alloy, specifically alloy 12, having a composition of Ti-48Al as is given in a number of the tables above. The one difference as noted int eh footnote to the table is that the binary TiAl alloy was prepared by cast and HIP processing rather than by the melt spinning and consolidation processing as set out in Examples 1-3 above.
Example 27 is an alloy similar to alloy 12 of Example 2b in that it contains the binary alloy but in this case the binary alloy is doped with 0.1 atom percent of boron. The processing of alloy 227 of Example 27 is essentially the same as the processing of alloy 12 of Example 2B and as is evident from a review of the data obtained by measuring yield strength, plastic elongation for samples annealed at temperatures ranging from 1250.degree. to 1350.degree. C., there is essentially no significant difference between the properties of the binary alloy of Example 2B and the doped binary 227 alloy of Example 27.
Considering next the alloy 133 of Example 26, this alloy contains 2 atom percent of chromium and 4 atom percent of niobium and is in this sense closely comparable to alloy 225 of Example 28 and alloy 246 of Example 29. Both of the latter alloys contain a boron dopant as well as the 2 atom percent of chromium and 4 atom percent of niobium. Each of these alloys, that is alloy 133, 225, and 246, was prepared by the cast and HIP processing as described above. If a comparison is made between the properties measured in tests of the respective alloys, it will be observed first that the yield strength of the undoped alloy 133 is relatively low and that the boron doped alloy 225 has a higher yield strength by only a relatively small measure. Similarly, the alloy 246 doped with 0.2 atom percent boron has a relatively low yield strength which is closely comparable to that of alloy 225 doped with 0.1 atom percent boron so that the level of doping of the two alloys with boron does not impart any significant change in strength. Further, there is very modest gain in strength over the alloy 133 which does not contain a boron dopant.
With regard next to the fracture strength, here again a modest increase in fracture strength is observed for the alloy 225 containing 0.1 atom percent boron dopant when compared with the alloy 133 which does not contain this dopant. Further, alloy 246 which contains 0.2 atom percent boron dopant does not have an increase in strength over the alloy 225 having 0.1 atom percent boron but rather has a modest decrease in strength.
With regard to the plastic elongation property for these three alloys, 133, 225, and 246, there does not appear to be a beneficial effect of the presence of the boron dopant in either the 0.1 atom percent or the 0.2 atom percent as compared to the same composition of alloy 133 which is free of the boron dopant.
EXAMPLES 26A through 29A
A number of additional samples were prepared by a cast and forged procedure as contrasted with the cast and HIP procedure of the examples 26 through 29 of Table VII. The chemistry of each of the alloys is essentially the same as that of the samples of Table VII. The difference between the samples is, accordingly, the difference in the method of preparation. The method of cast and forge processing is essentially as described above with reference to Example 24.
The specific alloy compositions homogenization temperatures, annealing temperatures, and physical properties of the alloys measured by tensile testing are listed in Table VIII immediately below.
In preparation of the samples of Table VIII, it will be noted that three of them were homogenized at 1250.degree. C. and that two, specifically 27A and 28A, were homogenized at 1400.degree. C.
A comparison of the data of the samples of Table VIII with the samples of Table VII reveal some important results. The ductility of the alloy 12 of Example 2A is considerably better than the ductibility of the same alloy is Example 2B of Table VII. The strength of the 2B alloy is essentially the same as that of the 2A alloy of Table VIII but there is an appreciable increase int he ductility of the samples prepared by the cast and forge processing when contrasted with the samples prepared by the cast and HIP processing of Table VII.
Alloy 227 of Example 27A is the binary alloy similar to that of Example 27 of Table VII and contains 0.1 atom percent boron. Alloy 227 of Example 27A was homogenized at 1400.degree. C. as contrasted with Example 27 of Table VII. Also, in Example 27A, the alloy was cast and forged as contrasted with the cast and HIP processing of Table VII. Considering the data listed for Example 27A in Table VIII in comparison with that for Example 27 of Table VII, it is evident that there is a gain in strength but there is also a reduction in ductility.
The incorporation of 0.1 atom percent boron in the alloy 225 of Example 28A does yield significant increase in ductility and this is evident from comparison of the data listed for Example 28A with the data listed for Example 26A. As is evident from Table VIII, two of the ductility values are over three and one is at a 3.5 level. This is an unusually high ductility for titanium aluminide. The significance of this data is that the combination of the doping with 0.1 atom percent boron and the homogenization treatment at 1400.degree. C. does yield significant improvement over the alloy 133 of Example 26A which contains no boron additive and which was homogenized at 1250.degree. C. It is also evident that the ductility values for Example 28A of Table VIII are far superior to the ductility values for the same sample, that is alloy 225, prepared according to the cast and HIP processing of Table VII. The conclusion is that the cast and forge processing and the higher temperature homogenization together with the boron doping does yield a ductility advantage which is evident by the comparisons described above with reference to Example 26A of Table VIII and with reference to Example 28 of Table VII.
The processing of the alloy 246 doped with 0.2 atom percent boron and homogenized at 1250.degree. C. does not yield significant advantage over the other alloys of Table VIII.
Accordingly, based on the foregoing, it is evident that a process for cast and forge preparation of alloys coupled with higher temperature homogenization and coupled also with boron doping does permit preparation of alloys having significantly higher ductility than is available from other processing procedures.
The increase in ductility possible by carrying out the procedure of the present invention is evident from FIG. 1 where the ductility data is plotted for the Example 26A compared to Example 28A.
What is provided pursuant to the present invention is a cast and wrought body of alloy. The alloy consists essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression:
The body is first cast and is then homogenized at a temperature close to or above the alpha transus temperature. By close to, as used herein, is meant within about thirty degrees of the transus temperature. The transus temperature is, of course, different for each alloy composition which falls within the above expression. Following the homogenization the body is forged to accomplish a deformation of at least ten percent. The combination of the chemistry of the alloy coupled with the high temperature homogenization and the forging imparts to the cast body the combination of desirable properties which are discussed above and illustrated in the table.
Claims
What is claimed is:
1. A cast and wrought body of alloy, said alloy consisting essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression: said body having been homogenized for one to three hours at a temperature close to or above the alpha transus temperature, and said body having been wrought to cause a deformation thereof of at least 10% and annealed.
2. A cast and wrought body of alloy, said alloy consisting essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression: said body having been homogenized for one to three hours at a temperature close to or above the alpha transus temperature, and said body having been wrought to cause a deformation thereof of at least 10% and annealed.
3. A cast and wrought body of alloy, said alloy consisting essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression: said body having been homogenized for one to three hours at a temperature close to or above the alpha transus temperature, and said body having been wrought to cause a deformation thereof of at least 10% and annealed.
4. A cast and wrought body of alloy, said alloy consisting essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression: said body having been homogenized for one to three hours at a temperature close to or above the alpha transus temperature, and said body having been wrought to cause a deformation thereof of at least 10% and annealed.
5. A cast and wrought body of alloy, said alloy consisting essentially of a gamma titanium aluminide modified by chromium, niobium, and boron according to the expression: said body having been homogenized for one to three hours at a temperature close to or above the alpha transus temperature, and said body having been wrought to cause a deformation thereof of at least 10% and annealed.
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