COBALT-BASE ALLOYS
United States Patent 3767385
Work-hardened cobalt-base alloys having high strength and ductility combined with excellent corrosion resistance, characterized by an iron content from 6 to 25 percent by weight but free of embrittling phases often associated with high iron content, the percent by weight of the other principal alloy constituents being: nickel, 18 - 40; molybdenum, 6 - 12; and chromium, 15 - 25. Further alloy embodiments comprising titanium and aluminum or columbium in a total amount up to 8 percent by weight and iron in an amount from 6 to 10 percent retain ductility at elevated temperatures to 900° F.
US Patent References:
Cobalt-nickel base alloys containing chromium and molybdenum
Smith - December 1967 - 3356542
Cobalt-nickel base alloys containing chromium and molybdenum
Smith - December 1967 - 3356542
Application Number:
05/174497
Publication Date:
10/23/1973
Filing Date:
08/24/1971
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Export Citation:
Assignee:
Standard Pressed Steel Co. (Jenkintown, PA)
Latrobe Steel Company (Latrobe, PA)
Latrobe Steel Company (Latrobe, PA)
Primary Class:
Other Classes:
419/28, 148/442, 148/557, 148/501
International Classes:
C22C19/07; C22F1/10; C22C19/00
Field of Search:
75/171,170,122,134F 148/32,32.5,31,11.5F,12.7
Primary Examiner:
Dean, Richard O.
Parent Case Data:
This application is a continuation-in-part application of copending application Ser. No. 76293 filed Sept. 28, 1970 (now abandoned).
Claims:
I claim
1. A cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases and consisting essentially of the following elements in percent by weight:
2. A cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases and consisting essentially of the following elements in percent by weight:
3. The method of making a cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases, which comprises forming a melt consisting essentially of the following elements in percent by weight:
4. The method of making a cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases, which comprises forming a melt consisting essentially of the following elements in percent by weight:
5. A cobalt-base alloy consisting essentially of the following elements in percent by weight:
6. An alloy as in claim 1 which is homogenized by heating prior to cold-working.
7. An alloy as in claim 1 which is additionally heat-aged after cold working.
8. An alloy as in claim 1 which is free of titanium, aluminum, and columbium.
9. An alloy as in claim 1 in which titanium is present together with at least aluminum or columbium.
10. An alloy as in claim 2 which has been cold-worked at ambient temperatures.
11. The method as in claim 3 wherein said alloy is homogenized by heating prior to cold-working.
12. The method as in claim 3 wherein said alloy is heat-aged after cold-working.
13. The method as in claim 3 wherein said alloy is powdered prior to cold working.
14. The method as in claim 13 wherein said powdered alloy is formed into a desired shape by compression and sintering prior to cold working.
15. The method as in claim 13 wherein said powdered alloy, prior to cold working, is sprayed hot onto a substrate to which it adheres.
16. The method as in claim 4 wherein said alloy is cold worked at ambient temperature.
1. A cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases and consisting essentially of the following elements in percent by weight:
2. A cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases and consisting essentially of the following elements in percent by weight:
3. The method of making a cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases, which comprises forming a melt consisting essentially of the following elements in percent by weight:
4. The method of making a cobalt-base alloy having a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase, said alloy being substantially free of embrittling phases, which comprises forming a melt consisting essentially of the following elements in percent by weight:
5. A cobalt-base alloy consisting essentially of the following elements in percent by weight:
6. An alloy as in claim 1 which is homogenized by heating prior to cold-working.
7. An alloy as in claim 1 which is additionally heat-aged after cold working.
8. An alloy as in claim 1 which is free of titanium, aluminum, and columbium.
9. An alloy as in claim 1 in which titanium is present together with at least aluminum or columbium.
10. An alloy as in claim 2 which has been cold-worked at ambient temperatures.
11. The method as in claim 3 wherein said alloy is homogenized by heating prior to cold-working.
12. The method as in claim 3 wherein said alloy is heat-aged after cold-working.
13. The method as in claim 3 wherein said alloy is powdered prior to cold working.
14. The method as in claim 13 wherein said powdered alloy is formed into a desired shape by compression and sintering prior to cold working.
15. The method as in claim 13 wherein said powdered alloy, prior to cold working, is sprayed hot onto a substrate to which it adheres.
16. The method as in claim 4 wherein said alloy is cold worked at ambient temperature.
Description:
The present invention relates to cobalt-base alloys having high strength and ductility combined with excellent corrosion resistance, said alloys comprising a hexagonal close-packed platelet phase dispersed in a face-centered cubic matrix phase and being substantially free of embrittling phases, and to methods of making such alloys.
U. S. Pat. No. 3,356,542 granted Dec. 5, 1967 to G. Smith discloses cobalt-nickel base alloys containing chromium and molybdenum. The alloys are corrosion-resistant and can be work-strengthened under certain temperature conditions to have very high ultimate tensile and yield strengths. These alloys can exist in one of two crystalline phases, depending on temperature. They are also characterized by a composition-dependent transition zone of temperatures in which transformations between the phases occur. At temperatures above the upper temperature limit of the transition zone, the alloys is stable in the face-centered cubic (FCC) structure. At temperatures below the lower temperature of the transformation zone, the alloys is stable in the hexagonal close-packed (HCP) form. By cold working metastable face-centered cubic material at a temperature below the lower limit of the transformation zone, some of it is transformed into the hexagonal close-packed phase which is dispersed as platelets throughout a matrix of the face-centered cubic material. It is this cold working and phase-transformation which is responsible for the excellent ultimate tensile and yield strength of the prior art alloys.
While the properties of these prior art alloys are desirable in many respects, particularly at room temperature, the alloys are quite expensive because of their high content of components such as nickel, molybdenum, and cobalt. The prior art materials are, in contrast, relatively poor in alloy components of low cost, such as iron. Iron may be present in the prior art alloys in amounts no greater than 6 percent by weight, for example. Attempts to increase the amount of low cost components such as iron in the alloys of the prior art have heretofore been discouraged because of the formation of topologically close-packed phases such as the sigma, mu, or chi phases (depending on composition) which severely embrittle the alloy.
Also, it has been found that those prior art alloys of highest ultimate tensile strength lose their ductility at relatively low temperatures, e.g. 300° - 400° F. This limits the field of use of these alloys to exclude high-temperature applications, such as in jet aircraft engines, or requires that a compromise be struck between the desired ductility, maximum service temperature, and ultimate tensile strength.
According to the present invention, highly corrosion resistant cobalt-based alloys having the excellent ultimate tensile and yield strengths imparted by formation of a platelet HCP phase in a matrix FCC phase have been prepared containing amounts of iron from 6 percent up to 23 percent while avoiding formation of embrittling phases. The resulting alloys have properties as good as, or nearly as good as, those of the prior art alloys but can be prepared at a lower cost to be more competitive with cheaper high-strength corrosion-resistant alloys. Further, certain of the alloys of the present invention which have high ultimate tensile strength also retain their ductility at elevated temperatures, e.g. at temperatures as high as 900° F.
In particular, the alloys of the present invention comprise the following chemical elements in the weight percent ranges indicated:
General Preferred Carbon 0 - 0.05 0 - 0.03 Nickel 18 - 40 18 - 30 Molybdenum 6 - 12 6 - 12 Chromium 15 - 25 18 - 22 Iron 6 - 25 or 7 - 23 or 6 - 10 7 - 10 Boron 0 - 0.1 0 - 0.03 Titanium 0 or 1 - 5 2 - 4 Aluminum 0 or 0 - 1 0.1 - 0.7 Columbium 0 or 0-2 0.1 - 1 Cobalt Balance Balance
As shown above, the alloys may or may not contain titanium. However, if the alloys are free of titanium, then they are also free of aluminum and columbium. On the other hand, if titanium is present, it is present in an amount from 1 - 5 weight percent, and either aluminum or columbium or both are present in amounts up to the maximum amounts indicated, e.g. if Ti = 0, Al + Cb = 0, and if Ti = 1 - 5, Al + Cb ≠ 0. Further, if titanium is present (with aluminum and/or columbium), the iron content of the alloys does not exceed 10 weight percent. As will be discussed more in detail hereinafter, those alloys containing titanium form a special class of materials combining high ultimate tensile strength with good ductility at elevated temperatures.
However, not all those alloys whose compositions fall within the ranges given above are encompassed by the present invention, since many of such compositions would include alloys containing embrittling phases. According to the present invention, it has been further found that the electron vacancy number, N v , defined as
N v = 0.61 Ni + 1.71 Co + 2.66 Fe + 4.66 Cr + 5.66 Mo ,
where the chemical symbols represent the effective atomic fractions (as defined later herein) of the respective elements present in the alloy, shall not exceed certain fixed values for alloys as described just above, in order to avoid the formation of embrittling phases. In particular, for those alloys containing from 16 to 23 or 25 percent of iron, the electron vacancy number N v should not exceed a value of 2.55. For alloys containing between 6 and 16 percent of iron, higher N v 's are permissible up to a maximum N v of 2.72 in alloys containing 6 percent by weight of iron. In those alloys containing between 6 and 16 percent of iron, the maximum permissible N v is defined as
N v = 2.82 - 0.017 W Fe where W Fe is the percent by weight of iron in the alloy. (Small effects on the N v due to carbon and boron are not significant in these alloys and may be ignored in these calculations.)
The N v number is a measure of the average number of electron vacancies in the 3d sub-band of chromium, cobalt, iron, and nickel, and the 4d sub-band of molybdenum. The formation of brittle phases in certain alloys can be correlated with these electron vacancies because of the involvement of these electrons in the chemical bonding of transition elements. An understanding of the definition and significance of the "effective atomic percentages" used in calculating the N v requires some further consideration of the microstructure of the alloys and of the correlation of structure with alloy properties.
As already mentioned, the conversion of a metastable FCC phase into platelets of stable HCP phase by cold working is believed to be responsible for the high tensile strength of the alloys of the prior art and of the present invention because, in theory, the formation of these platelets of the HCP phase in the FCC matrix inhibits the relative movement along slip planes which is responsible for failure in metals. However, ductility in these alloys can also be correlated with the phenomenon of phase conversion and, indeed, the surprising ductility of these alloys has been attributed to the occurrence of phase conversion during deformation such as in a tensile test. In any case, it is clear that the greater is the degree of cold working and the higher is the ultimate tensile strength in these alloys, the lower their ductility becomes.
Thus, the alloys, when cold worked to increase their strength, lose ductility. Further, the remaining ductility decreases as the temperature rises. The high-temperature loss of ductility in the alloys, for example in a fastener used in a jet aircraft engine, reduces the ability of the fastener to accommodate to a local concentration of stress, for example, and may result in failure, whereas a more ductile material would deform to accommodate and equalize such stress.
One way to retain ductility in the prior art alloys is to decrease the degree to which they are cold worked. However, the less-worked alloys will also have a lower ultimate tensile strength. Those alloys of the present invention which contain titanium with aluminum and/or columbium resolve this dilemma: a high ultimate tensile strength is produced with a lower degree of cold working (and greater preservation of ductility at elevated tmperatures) than in those alloys free of these elements.
It is believed that these elements form compounds of the formula Ni 3 X, where X is titanium, aluminum, and/or columbium. (Some substitution of other elements in such compounds is believed to occur, as discussed more in detail below.) It is believed further that the formation of these compounds results in a precipitation hardening of the alloys in question, supplementing the hardening effect due to the conversion of FCC to HCP phase, and thus permitting the attainment of a higher tensile strength at a lower degree of cold working than heretofore possible.
The "effective atomic fraction" of elements in these alloys, as used to calculate the electron vacancy number, N v , takes into account the postulated conversion of a portion of the metal atoms present, particularly nickel, into compounds of the type Ni 3 X.
Thus, for those alloys which contain no titanium (and hence no aluminum or columbium), the "effective atomic fraction" of each element present is the same as its actual atomic fraction. No Ni 3 X compounds are formed in these alloys.
For alloys containing the "X" or "hardener" elements (titanium, aluminum,and/or columbium), the total atomic percent of each of the elements present in a given alloy is first calculated from the weight percent ignoring the carbon and/or boron in the composition. Each atomic percentage represents the number of atoms of an element present in 100 atoms of alloy. The number of atoms/100 (or atomic percentage) of hardener elements is totalled and multiplied by 4 to give an approximate number of atoms/100 involved in Ni 3 X formation. This figure, however, must be adjusted.
R. W. Guard et al., in"The Alloying Behavior of Ni 3 Al (Gamma-Prime Phase)," Met. Soc. AIME 215, 807 (1959), have shown that cobalt, iron, chromium, and molybdenum enter such an Ni 3 X compound in amounts up to 23, 15, 16,and 1 percent, respectively. To approximate the number of atoms/100 of each of these metals which are also "tied up" in the Ni 3 X phase and are unavailable for formation of non-Ni 3 X matrix alloy, the product of the maximum percent solubility of each metal in Ni 3 X, its atomic fraction in the alloy under consideration, and the total number of atoms of Ni 3 X possible in 100 atoms of alloy is found.
The number of atoms of Ni, Co, Fe, Cr, and Mo in 100 atoms of alloy, respectively, are then corrected by subtraction of the figures representing the amount of each of these metals in the Ni 3 X phase. The difference approximates the number of atoms per 100 of the nominal alloy composition which are effectively available for matrix alloy formation. Since this total number is less than 100, the "effective atomic percent" of each of the elements -- based on this total -- is now calculated. The effective atomic fraction, which is the quotient of the effective atomic percent divided by 100, is employed in the determination of N v for these alloys. This calculation is exemplified in detail later herein.
For preparing the alloys of the present invention, the components are first melted, suitably by vacuum induction melting, at a temperature of about 2,700° F. Because of the relatively high content of iron in the alloys, ferroalloys of chromium and molybdenum can be used in preparing the melt, leading to a considerable cost advantage. Typically, most of the iron would be added in the form of ferroalloys.
The molten alloys may be cast as ingots, or can be impinged by a gas jet or on a surface to disperse the melt as small droplets. The latter solidify to form a metallic powder useful in powder metallurgy processes hereinafter described.
If desired, ingots cast from the melt prepared as above can be further processed, after cooling, in a vacuum consumable electrode furnace to improve the ingot structure and to reduce the gas content.
It is well-known to those skilled in the art that micro-segregation, i.e. "coring" occurs during the solidification of alloys, particularly in ingot form, under practical working conditions. This micro-segregation results in the formation of dendrites in the micro-structures and, in the alloys of the present invention, has the effect of increasing the N v in the material between the dendrites. This is because the material first to solidify as the alloy melt cools has a relatively high content of cobalt and nickel, which have little tendency to form embrittling phases such as the sigma phase. In contrast, the last solidifying material, which is relatively rich in chromium and molybdenum, has a higher electron vacancy number and tends to form embrittling phases. Accordingly, certain alloy compositions according to the present invention, if of a composition tending toward the maximum permissible N v , may contain sigma phase in the as-cast condition. Thus, the alloys of the present invention are preferably homogenized by heating at temperatures below their melting point, suitably at temperatures from about 2,050° F. to about 2,150° F., for 18 - 36 hours to diffuse the elements and so eliminate any sigma or other embrittling phases which may possibly be present in the alloy as cast.
After homogenization, alloy ingots are usually pressed preparatory to rolling and then are rolled to shapes and sizes convenient for further working. (Alternatively, the ingots may be reduced to powdered form after homogenization and prior to any working.) Pressing, if necessary for large ingots, is carried out at temperatures of from 2,000° F. to 2,100° F., and may bring about a reduction in area of about 4 to 1. Rolling is carried out at the same temperature and may cause a further reduction in cross-section of about 40 to 1.
The rolled material may optionally be subsequently annealed at a temperature between about 1,925° - 1,950° F. for about 1 hour to achieve a uniform structure for subsequent cold working.
The cold working takes place at a temperature below the lower temperature of the temperature zone of transition from the high-temperature face-centered cubic phase to the low-temperature stable hexagonal close-packed phase. Generally, cold working is most conveniently effected at ambient temperatures, which may vary in a conventional mill from about 0° F. to 110° F., for example. These ambient temperatures are well below the lower temperature of the transition zone for all alloys encompassed by the present invention.
In case cold working at a temperature above ambient temperature is desired, the temperature limits of the transformation zone can be quite simply determined for any particular alloy composition. For example, a sample of the alloy is first cold worked at ambient temperatures to give a reduction in cross-section of about 50 percent. Cold working under these conditions brings about phase transformation resulting in the formation of a platelet HCP phase in a matrix FCC phase. These phase differences are clearly evident on inspection of electron micrographs (magnification of 2,000 - 3,000 diameters) of the alloy. Different portions of the cold-worked sample, such as slices from a rod or bar, are then heated for about 8 hours each under temperature conditions differing incrementally by 50° or 100° F. over a range of temperatures between 800° F. and 1,500° F., depending on the alloy composition. Electron micrographs of each differently heat-treated portion are then examined. The temperature at which the HCP platelet phase begins to disappear from the original two-phased structure is the lower temperature limit of the transformation zone. The temperature at which all of the HCP platelet phase has disappeared, leaving only the FCC matrix phase, is the upper temperature limit of the transformation zone for an alloy of the particular composition studied. For the alloy of Example 1 below, for instance, the transformation zone lies approximately between 1,100° F. and 1,350° F.
Cold working of cast alloys at a temperature below the lower limit of the transformation zone may be by conventional practical means such as rolling, swaging, or drawing. Cold drawing or rolling are preferred, however, because they produce a more uniform product. When drawing to round shapes, the annealed materials are first submitted to centerless grinding to remove hot-rolling scale and to obtain a uniform bar size so that further area reduction by cold-working can be more readily controlled. After centerless grinding, the stock is suitably lubricated, for example with a coating of lime and/or molybdenum disulfide, and then cold-drawn.
Powdered alloys, prepared as described earlier herein may be handled in different ways. For instance, the powders may be hot- or cold-pressed into a desired shape and then sintered according to techniques known in powder metallurgy. The shaped articles can be cold worked by coining, for example. In another technique, the powdered alloy is sprayed hot (e.g. by so-called "plasma spraying") onto a substrate, such as of another metal, to which it adheres, and then cold worked in situ by suitable means such as swaging, rolling. or hammering.
Cold working is carried out to bring about a reduction in cross-section of at least about 5 percent. The most useful tensile strengths develop in alloys which have been cold worked to effect a reduction in cross-section of at least about 30 percent. As discussed earlier, the more the alloys are cold-worked, the more strength they develop. However, as the alloys increase in strength, they become less ductile. Accordingly, although the alloys can be worked to achieve a reduction in cross-section of as much as 95 percent, in practice they are not worked to an area reduction greater than 70 percent when the alloy is used, for example, for the manufacture of wire. Where still greater ductility is desired, as for example in alloys which are to be manufactured into fasteners, a reduction in cross-section between 50 - 60 percent is generally sufficient. In those alloys accordng to the invention which contain titanium with aluminum and/or columbium, desired high ultimate tensile strengths (260 - 280 ksi) can be obtained with a reduction in cross-section between about 35 and 45 percent. These alloys have particularly good ductility and retain it even at elevated temperatures.
To increase the strength of the cold worked alloys, the alloys may next be aged at a temperature between 800° F. and 1,350° F. for about 4 hours. A preferred aging temperature for those alloys free of titanium is about 900° F. If titanium is present, higher aging temperatures of about 1,200° F. are preferred. After aging, the materials are air-cooled.
It is pointed out that the alloys disclosed herein may also be worked at temperatures above the upper limit of the transformation zone (so-called "hot working") to produce articles of a desired shape. Such alloys, however, will not have the same high tensile strength characteristic of those alloys which are cold worked as hereinbefore described.
Also, it is possible to hot work those alloys of the present invention which have been priorly cold worked. Often, whatever loss of tensile strength is brought about by hot working is acceptable. Thus, for example, articles can be made by the partial hot forming (e.g. at about 1,900° F.) of previously cold worked stock. Although there is some loss in tensile strength in the hot worked portion, this may be acceptable if the hot worked portion is subjected to lower stresses than are those portions which have not been hot worked. The resultant article thus can combine desired strength properties with convenient and relatively low cost manufacture.
A better understanding of the present invention and of its many advantages will be had by referring to the following specific Examples, given by way of illustration.
EXAMPLE 1
An alloy, designated J80, having the following composition:
Percent by Weight C 0.01 Mo 9.9 Cr 19.4 Co 30.4 Fe 17.3 Ni 23.0 B 0.01
and having an electron vacany number, N v , of 2.52 was prepared from an approximately 50 pound charge comprising 9.90 lb of electrolytic chromium, 11.52 lb of nickel, 15.30 lb of cobalt, 5.10 lb of molybdenum, 0.04 lb of nickel-boron, and 8.70 lb of electrolytic iron. These components were melted in a MgO crucible in a vacuum furnace under 10 microns of mercury pressure at about 2,700° F. The alloy was refined until the leak-up rate, measured at 5-minute intervals, was substantially constant.
The melt was poured into a cylindrical steel mold with a copper stool, and then solidified. After cooling, the vacuum was broken, the furnace opened, and the mold stripped from the ingot, which was about 3.5 inches in diameter and 15 inches long.
Because the shape of the ingot permitted, the ingot was directly rolled at 2,050° F. to reduce its diameter to 0.75 inch, or an area reduction of about 22 to 1.
The rolled ingot was next annealed for 1 hour at 1,925° F., then centerless ground to a diameter of about 0.71 inch.
The ground piece was next drawn to 0.5 inch diameter using a lubricant of lime and molybdenum disulfide. The resulting 50 percent reduction in cross-section was achieved on about three passes, each effecting a 20 percent reduction in cross-section. The cold working was effected at ambient temperature, i.e. about 70° F.
The cold worked alloy was next aged for 4 hours at 950° F. and air-cooled. Specimens were then cut and ground for testing. The alloy showed the following properties. (ksi = kilopounds per square inch = 1,000 psi)
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 278.6 271.5 6.5 31.3 274.4 268.4 6.5 30.4
If the same alloy is cold worked to only a 38 percent reduction in cross section at 70° F., the following properties are found:
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 217.2 200.4 7.5 18.4 217.1 200.8 7.5 32.4
Although the alloys of the present invention contain significantly more iron than those taught in the aforementioned patent U. S. Pat. No. 3,356,542 to Smith, they surprisingly have a greater corrosion resistance than the prior art alloys.
In one test of corrosion resistance, a controlled voltage is applied by a potentiostat across a cell having a saturated calomel cathode and the alloy to be tested as the anode. A 3.5 percent NaCl solution is used as the electrolyte. The current through the cell is measured: by Coulomb's Law, the smaller the current observed, the smaller is the corrosion.
The J80 alloy of this Example, annealed and cold worked to an area reduction of 50 percent, showed the following corrosion characteristics in such a test:
Voltage Current (milliamperes) 0.2 nil 0.4 0.015 0.6 0.05 0.8 0.095
(Open circuit potential v. saturated calomel = 0.3 v.).
In contrast, a typical prior art alloy free of iron (whose composition is given below) showed the following results:
Voltage Current (milliamperes) 0.2 0.004 0.4 0.02 0.6 0.13 0.8 1.2
(Open circuit potential v. saturated calomel = 0.174 v.) This prior art alloy, which had been annealed and cold worked to an area reduction of 47.5 percent, had the following composition:
Percent by Weight C 0.02 Mo 10 Cr 20 Co 35 Ni 35
In another test for measuring the resistance of the alloys to crevice corrosion (i.e. corrosion under anaerobic conditions in which an oxygen concentration gradient is established, resulting in the formation of a concentration cell), bars of the J80 and prior art alloy were wrapped longitudinally with a stretched rubber band and immersed in 10 percent aqueous FeCl 3 solution for 72 hours at room temperature. No crevice corrosion was observed in any of the samples under these conditions. The degree of general surface corrosion was measured by determining weight loss per unit surface area.
Weight Loss (mg/in 2 ) J80 (cold worked 38 percent) 1.1 J80 (cold worked 50 percent) 1.2 Prior Art (cold worked 30 percent) 1.4 Prior Art (cold worked 30 percent) 3.2
(All four alloys were heat-aged for four hours at 1,000° F. after cold-working). Despite the high iron content of the alloy of the present invention, these tests show that both the crevice corrosion resistance and the general corrosion resistance of the alloy are, unexpectedly, comparable with the same properties of an iron-free prior art nickel-cobalt alloy.
EXAMPLES 2 - 7
A number of further alloy bodies were prepared as in Example 1 above having the composition shown below in Table I. While the percent by weight composition of each of the alloys falls within the ranges disclosed earlier herein, the N v value for certain of the alloys falls outside the permissible limits disclosed. Such alloys are not within the scope of the present invention and are shown for purposes of comparison. (All alloys contained 0.01 percent boron and 0.015 percent carbon in addition to the elements listed.)
TABLE I
Mo Cr Co Fe Ni N v J 83 10.0 20.0 31.5 9.5 29.0 2.40 J84 7.0 20.0 40.0 12.0 21.0 2.43 J88 8.5 20.2 30.3 19.3 21.7 2.54 J81 10.4 19.8 31.6 17.0 21.2 2.56 J82 10.8 20.3 32.8 16.7 19.4 2.61 J85 9.0 20.5 31.5 19.0 20.0 2.58
The alloys J81, J82, and J85 all contained a heavy sigma phase and were too brittle to be processed. Average properties for the remaining alloys after processing as in Example 1 above are reported in Table II.
TABLE II
Ultimate Tensile 0.2 Yield Reduction in Strength Strength Elongation Area Alloy (ksi) (ksi) (%) (%) J83 289.3 274.7 2.7 8.6 J84 287.7 256.9 5.2 21.9 J88 261.4 241.8 8.2 40.4
example 8
when preparing alloy bodies on a large scale, it is convenient to use ferroalloys in the melt. Thus, a 1,500 pound charge was prepared by vacuum melting the following components: 269 lb of low-carbon ferrochrome; 148 lb of pure molybdenum; 348 lb of nickel; 456 lb of cobalt; 88 lb of chromium; 190 lb of iron; 1 lb of nickel-boron; and 0.1 lb of additional carbon. A wet analysis of the alloy was as follows:
C 0.011 Cr 18.96 Si 0.01 Mo 0.86 Mn 0.03 Co 30.33 S 0.003 Fe 17.60 P 0.008 B 0.01 Ni Balance
EXAMPLE 9
The beneficial effects of heat aging after cold working are shown by comparing the properties earlier reported for the cold-drawn and heat-aged alloy of Example 1 with the properties (given below) for the same alloy after cold-working, but prior to heat-aging.
50 Percent Drawn (Average of Two Samples)
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 218 205 8.8 41.0
38 Percent Drawn (Average of Two Samples)
195 150 13.5 45.5 EXAMPLE 10 Proceeding as in Example 1, the following alloy, designated as J153, was prepared by melting, casting, and rolling: Percent by Weight C 0.015 Mo 7.0 Cr 19.0 Co 35.7 Fe 9.0 Ni 25.5 B 0.015 Ti 3.0 Al 0.2 Cb 0.6
The effective atomic percent and N v for this alloy are calculated as follows.
The atom percent of each element present is first calculated, ignoring boron and carbon:
Atom Percent (No. of atoms per 100) Mo 4.25 Cr 21.29 Co 35.30 Fe 9.39 Ni 25.31 Ti 3.65 Al 0.43 Cb 0.38 Total 100.00
The number of hardener element ("X") atoms/100 alloy atoms is (3.65 + 0.43 + 0.38) = 4.46. The total unadjusted number of nickel and X atoms involved in Ni 3 X formation is (4 × 4.46) = 17.84, of which 13.38 are nickel, leaving (25.31 - 13.38) = 11.93 atoms of nickel available for matrix formation.
The number of molybdenum atoms which enter the Ni 3 X phase is the product of the Mo solubility in Ni 3 X, the atomic fraction of Mo in the alloy, and the number of Ni 3 X atoms, or (0.01 × 0.0425 × 17.84) = 0.01, leaving (4.25 - 0.01) = 4.24 atoms of Mo per 100 of alloy composition for matrix formation.
For chromium, (0.16 × 0.2129 × 17.84) = 0.61. Available Cr = (21.29 - 0.61) = 20.68 atoms/100.
For cobalt, (0.23 × 0.3530 × 17.84) = 1.44. Available Co = (35.30 - 1.44) = 33.86 atoms/100.
For iron, (0.15 × 0.0939 × 17.84) = 0.25. Available Fe = (9.39 - 0.25) = 9.14 atoms/100.
The total number of atoms, the number of atoms available for matrix formation, the effective atom percent, based on this latter total, and the effective atom fraction, are as follows:
No./100 Available Effective Effective (Atom %) No./100 Atom Percent atom Fraction Mo 4.25 4.24 5.31 .0531 Cr 21.29 20.68 25.90 .2590 Co 35.30 33.86 42.39 .4239 Fe 9.39 9.14 11.45 .1145 Ni 25.31 11.93 14.95 .1495 Ti 3.65 Al 0.43 Cb 0.38 Total 100.00 79.85 100.00 1.00
The N v for this alloy is 2.63, calculated by substituting these effective atom fractions into the equation
N v = 0.61 Ni + 1.71 Co + 2.66 Fe + 4.66 Cr + 5.66 Mo .
This value is less than the maximum permissible N v of 2.67 for an alloy containing 9 weight percent of iron from
N v = -0.017 W Fe + 2.82.
The alloy, thus, is substantially free of embrittling phases.
The alloy was next cold worked to reduce its cross-section by 40 percent, and then aged for 4 hours at 1,225° F. Samples of the alloy were tested at room temperature, 700° F., and at 900° F. to determine their strength and ductility, as reported in Table III below. Notched tensile strength (notch factor, K T = 6) and ultimate tensile strength were measured at each temperature and are compared in the Table.
TABLE III
Test Ultimate 0.2% Yield Temp. Tensile Strength Reduction in (°F) Strength (ksi) Elongation Area (ksi) (%) (%) R.T. 274.3 264.6 9.5 44.4 R.T. 268.7 256.6 11.0 43.8 Total 271.5 260.6 10.3 44.1 R.T. 379.4 NTS:UTS = 1.4:1 R.T. 380.4 Total 379.9 700 225.1 218.4 10.0 37.7 700 227.7 220.4 9.0 34.5 Total 226.4 219.4 9.5 36.1 700 320.8 NTS:UTS = 1.44:1 700 324.5 Total 322.7 900 205.6 201.2 10.0 33.0 900 202.8 197.6 11.0 37.3 Total 204.2 199.4 10.5 35.2 900 311.1 NTS:UTS = 1.50:1 900 299.8 Total 305.5
The effects of cold working on tensile strength and ductility can be seen by a comparison of Table III with following Table IV showing the same alloy properties in samples of the alloy of Table III cold worked to reduce their cross section by 46 percent.
TABLE IV
Ultimate Test Tensile 0.2% Yield Reduction Temp Strength Strength Elongation in Area (°F) (ksi) (ksi) (%) (%) R.T. 288.9 272.7 9.0 39.6 R.T. 282.5 273.6 8.5 36.2 Total 285.7 273.2 8.8 37.8 R.T. 388.5 NTS:UTS = 1.35:1 386.1 Total 387.3 700 226.5 220.4 8.0 26.7 700 232.3 222.2 9.0 33.3 Total 229.4 221.3 8.5 30.0 700 329.8 NTS:UTS = 1.46:1 700 332.8 Total 331.3 900 226.6 220.7 9.5 31.8 900 224.4 218.8 9.5 33.3 Total 225.5 219.8 9.5 32.6 900 317.7 NTS:UTS = 1.40:1 900 315.1 Total 316.4
EXAMPLE 11
Another alloy, designated J150, of lower iron content than the alloy of Example 10, was prepared as in Example 1. The composition of the alloy is given below. The N v of the alloy is 2.64, or below the permissible maximum of 2.70. The physical properties of samples cold worked to reduce their cross section by 40 percent and then aged at 1,225° F. for 4 hours are reported at various temperatures in Table V.
ALLOY COMPOSITION
Percent by Weight C 0.015 Mo 7.0 Cr 20.0 Co 36.2 Fe 7.0 Ni 26.0 B 0.015 Ti 3.0 Al 0.2 Cb 0.6
TABLE V
Test Ultimate 0.2% Yield Reduction Temp Tensile Strength Elongation in Area (°F) Strength (ksi) (%) (%) (ksi) R.T. 291.4 284.4 10.5 39.1 R.T. 286.3 282.3 10.5 42.9 Total 288.9 283.4 10.5 41.0 R.T. 382.9 NTS:UTS = 1.34:1 R.T. 392.6 Total 387.8 700 241.4 236.6 10.0 34.4 700 241.5 240.3 9.0 33.4 Total 241.4 238.5 9.5 33.9 700 350.2 NTS:UTS = 1.43:1 700 339.9 Total 345.0 900 240.0 230.5 4.0 8.6 900 237.7 228.5 5.0 10.0 Total 238.9 229.5 4.5 9.3 900 333.1 NTS:UTS = 1.38:1 900 329.3 Total 331.2
EXAMPLE 12
Following Table VI summarizes the nominal chemical composition of three alloys. Alloy A is a prior art iron-free alloy falling within the scope of U. S. Pat. No. 3,356,542 to Smith. Alloy B is an alloy according to the present invention, like the alloy of Example 1. Alloy C is an alloy like that in Example 10.
Table VII below compares the physical properties of these alloys at various temperatures after cold working to the same approximate degree of reduction in cross section and aging for 4 hours. ##SPC1## ##SPC2##
U. S. Pat. No. 3,356,542 granted Dec. 5, 1967 to G. Smith discloses cobalt-nickel base alloys containing chromium and molybdenum. The alloys are corrosion-resistant and can be work-strengthened under certain temperature conditions to have very high ultimate tensile and yield strengths. These alloys can exist in one of two crystalline phases, depending on temperature. They are also characterized by a composition-dependent transition zone of temperatures in which transformations between the phases occur. At temperatures above the upper temperature limit of the transition zone, the alloys is stable in the face-centered cubic (FCC) structure. At temperatures below the lower temperature of the transformation zone, the alloys is stable in the hexagonal close-packed (HCP) form. By cold working metastable face-centered cubic material at a temperature below the lower limit of the transformation zone, some of it is transformed into the hexagonal close-packed phase which is dispersed as platelets throughout a matrix of the face-centered cubic material. It is this cold working and phase-transformation which is responsible for the excellent ultimate tensile and yield strength of the prior art alloys.
While the properties of these prior art alloys are desirable in many respects, particularly at room temperature, the alloys are quite expensive because of their high content of components such as nickel, molybdenum, and cobalt. The prior art materials are, in contrast, relatively poor in alloy components of low cost, such as iron. Iron may be present in the prior art alloys in amounts no greater than 6 percent by weight, for example. Attempts to increase the amount of low cost components such as iron in the alloys of the prior art have heretofore been discouraged because of the formation of topologically close-packed phases such as the sigma, mu, or chi phases (depending on composition) which severely embrittle the alloy.
Also, it has been found that those prior art alloys of highest ultimate tensile strength lose their ductility at relatively low temperatures, e.g. 300° - 400° F. This limits the field of use of these alloys to exclude high-temperature applications, such as in jet aircraft engines, or requires that a compromise be struck between the desired ductility, maximum service temperature, and ultimate tensile strength.
According to the present invention, highly corrosion resistant cobalt-based alloys having the excellent ultimate tensile and yield strengths imparted by formation of a platelet HCP phase in a matrix FCC phase have been prepared containing amounts of iron from 6 percent up to 23 percent while avoiding formation of embrittling phases. The resulting alloys have properties as good as, or nearly as good as, those of the prior art alloys but can be prepared at a lower cost to be more competitive with cheaper high-strength corrosion-resistant alloys. Further, certain of the alloys of the present invention which have high ultimate tensile strength also retain their ductility at elevated temperatures, e.g. at temperatures as high as 900° F.
In particular, the alloys of the present invention comprise the following chemical elements in the weight percent ranges indicated:
General Preferred Carbon 0 - 0.05 0 - 0.03 Nickel 18 - 40 18 - 30 Molybdenum 6 - 12 6 - 12 Chromium 15 - 25 18 - 22 Iron 6 - 25 or 7 - 23 or 6 - 10 7 - 10 Boron 0 - 0.1 0 - 0.03 Titanium 0 or 1 - 5 2 - 4 Aluminum 0 or 0 - 1 0.1 - 0.7 Columbium 0 or 0-2 0.1 - 1 Cobalt Balance Balance
As shown above, the alloys may or may not contain titanium. However, if the alloys are free of titanium, then they are also free of aluminum and columbium. On the other hand, if titanium is present, it is present in an amount from 1 - 5 weight percent, and either aluminum or columbium or both are present in amounts up to the maximum amounts indicated, e.g. if Ti = 0, Al + Cb = 0, and if Ti = 1 - 5, Al + Cb ≠ 0. Further, if titanium is present (with aluminum and/or columbium), the iron content of the alloys does not exceed 10 weight percent. As will be discussed more in detail hereinafter, those alloys containing titanium form a special class of materials combining high ultimate tensile strength with good ductility at elevated temperatures.
However, not all those alloys whose compositions fall within the ranges given above are encompassed by the present invention, since many of such compositions would include alloys containing embrittling phases. According to the present invention, it has been further found that the electron vacancy number, N v , defined as
N v = 0.61 Ni + 1.71 Co + 2.66 Fe + 4.66 Cr + 5.66 Mo ,
where the chemical symbols represent the effective atomic fractions (as defined later herein) of the respective elements present in the alloy, shall not exceed certain fixed values for alloys as described just above, in order to avoid the formation of embrittling phases. In particular, for those alloys containing from 16 to 23 or 25 percent of iron, the electron vacancy number N v should not exceed a value of 2.55. For alloys containing between 6 and 16 percent of iron, higher N v 's are permissible up to a maximum N v of 2.72 in alloys containing 6 percent by weight of iron. In those alloys containing between 6 and 16 percent of iron, the maximum permissible N v is defined as
N v = 2.82 - 0.017 W Fe where W Fe is the percent by weight of iron in the alloy. (Small effects on the N v due to carbon and boron are not significant in these alloys and may be ignored in these calculations.)
The N v number is a measure of the average number of electron vacancies in the 3d sub-band of chromium, cobalt, iron, and nickel, and the 4d sub-band of molybdenum. The formation of brittle phases in certain alloys can be correlated with these electron vacancies because of the involvement of these electrons in the chemical bonding of transition elements. An understanding of the definition and significance of the "effective atomic percentages" used in calculating the N v requires some further consideration of the microstructure of the alloys and of the correlation of structure with alloy properties.
As already mentioned, the conversion of a metastable FCC phase into platelets of stable HCP phase by cold working is believed to be responsible for the high tensile strength of the alloys of the prior art and of the present invention because, in theory, the formation of these platelets of the HCP phase in the FCC matrix inhibits the relative movement along slip planes which is responsible for failure in metals. However, ductility in these alloys can also be correlated with the phenomenon of phase conversion and, indeed, the surprising ductility of these alloys has been attributed to the occurrence of phase conversion during deformation such as in a tensile test. In any case, it is clear that the greater is the degree of cold working and the higher is the ultimate tensile strength in these alloys, the lower their ductility becomes.
Thus, the alloys, when cold worked to increase their strength, lose ductility. Further, the remaining ductility decreases as the temperature rises. The high-temperature loss of ductility in the alloys, for example in a fastener used in a jet aircraft engine, reduces the ability of the fastener to accommodate to a local concentration of stress, for example, and may result in failure, whereas a more ductile material would deform to accommodate and equalize such stress.
One way to retain ductility in the prior art alloys is to decrease the degree to which they are cold worked. However, the less-worked alloys will also have a lower ultimate tensile strength. Those alloys of the present invention which contain titanium with aluminum and/or columbium resolve this dilemma: a high ultimate tensile strength is produced with a lower degree of cold working (and greater preservation of ductility at elevated tmperatures) than in those alloys free of these elements.
It is believed that these elements form compounds of the formula Ni 3 X, where X is titanium, aluminum, and/or columbium. (Some substitution of other elements in such compounds is believed to occur, as discussed more in detail below.) It is believed further that the formation of these compounds results in a precipitation hardening of the alloys in question, supplementing the hardening effect due to the conversion of FCC to HCP phase, and thus permitting the attainment of a higher tensile strength at a lower degree of cold working than heretofore possible.
The "effective atomic fraction" of elements in these alloys, as used to calculate the electron vacancy number, N v , takes into account the postulated conversion of a portion of the metal atoms present, particularly nickel, into compounds of the type Ni 3 X.
Thus, for those alloys which contain no titanium (and hence no aluminum or columbium), the "effective atomic fraction" of each element present is the same as its actual atomic fraction. No Ni 3 X compounds are formed in these alloys.
For alloys containing the "X" or "hardener" elements (titanium, aluminum,and/or columbium), the total atomic percent of each of the elements present in a given alloy is first calculated from the weight percent ignoring the carbon and/or boron in the composition. Each atomic percentage represents the number of atoms of an element present in 100 atoms of alloy. The number of atoms/100 (or atomic percentage) of hardener elements is totalled and multiplied by 4 to give an approximate number of atoms/100 involved in Ni 3 X formation. This figure, however, must be adjusted.
R. W. Guard et al., in"The Alloying Behavior of Ni 3 Al (Gamma-Prime Phase)," Met. Soc. AIME 215, 807 (1959), have shown that cobalt, iron, chromium, and molybdenum enter such an Ni 3 X compound in amounts up to 23, 15, 16,and 1 percent, respectively. To approximate the number of atoms/100 of each of these metals which are also "tied up" in the Ni 3 X phase and are unavailable for formation of non-Ni 3 X matrix alloy, the product of the maximum percent solubility of each metal in Ni 3 X, its atomic fraction in the alloy under consideration, and the total number of atoms of Ni 3 X possible in 100 atoms of alloy is found.
The number of atoms of Ni, Co, Fe, Cr, and Mo in 100 atoms of alloy, respectively, are then corrected by subtraction of the figures representing the amount of each of these metals in the Ni 3 X phase. The difference approximates the number of atoms per 100 of the nominal alloy composition which are effectively available for matrix alloy formation. Since this total number is less than 100, the "effective atomic percent" of each of the elements -- based on this total -- is now calculated. The effective atomic fraction, which is the quotient of the effective atomic percent divided by 100, is employed in the determination of N v for these alloys. This calculation is exemplified in detail later herein.
For preparing the alloys of the present invention, the components are first melted, suitably by vacuum induction melting, at a temperature of about 2,700° F. Because of the relatively high content of iron in the alloys, ferroalloys of chromium and molybdenum can be used in preparing the melt, leading to a considerable cost advantage. Typically, most of the iron would be added in the form of ferroalloys.
The molten alloys may be cast as ingots, or can be impinged by a gas jet or on a surface to disperse the melt as small droplets. The latter solidify to form a metallic powder useful in powder metallurgy processes hereinafter described.
If desired, ingots cast from the melt prepared as above can be further processed, after cooling, in a vacuum consumable electrode furnace to improve the ingot structure and to reduce the gas content.
It is well-known to those skilled in the art that micro-segregation, i.e. "coring" occurs during the solidification of alloys, particularly in ingot form, under practical working conditions. This micro-segregation results in the formation of dendrites in the micro-structures and, in the alloys of the present invention, has the effect of increasing the N v in the material between the dendrites. This is because the material first to solidify as the alloy melt cools has a relatively high content of cobalt and nickel, which have little tendency to form embrittling phases such as the sigma phase. In contrast, the last solidifying material, which is relatively rich in chromium and molybdenum, has a higher electron vacancy number and tends to form embrittling phases. Accordingly, certain alloy compositions according to the present invention, if of a composition tending toward the maximum permissible N v , may contain sigma phase in the as-cast condition. Thus, the alloys of the present invention are preferably homogenized by heating at temperatures below their melting point, suitably at temperatures from about 2,050° F. to about 2,150° F., for 18 - 36 hours to diffuse the elements and so eliminate any sigma or other embrittling phases which may possibly be present in the alloy as cast.
After homogenization, alloy ingots are usually pressed preparatory to rolling and then are rolled to shapes and sizes convenient for further working. (Alternatively, the ingots may be reduced to powdered form after homogenization and prior to any working.) Pressing, if necessary for large ingots, is carried out at temperatures of from 2,000° F. to 2,100° F., and may bring about a reduction in area of about 4 to 1. Rolling is carried out at the same temperature and may cause a further reduction in cross-section of about 40 to 1.
The rolled material may optionally be subsequently annealed at a temperature between about 1,925° - 1,950° F. for about 1 hour to achieve a uniform structure for subsequent cold working.
The cold working takes place at a temperature below the lower temperature of the temperature zone of transition from the high-temperature face-centered cubic phase to the low-temperature stable hexagonal close-packed phase. Generally, cold working is most conveniently effected at ambient temperatures, which may vary in a conventional mill from about 0° F. to 110° F., for example. These ambient temperatures are well below the lower temperature of the transition zone for all alloys encompassed by the present invention.
In case cold working at a temperature above ambient temperature is desired, the temperature limits of the transformation zone can be quite simply determined for any particular alloy composition. For example, a sample of the alloy is first cold worked at ambient temperatures to give a reduction in cross-section of about 50 percent. Cold working under these conditions brings about phase transformation resulting in the formation of a platelet HCP phase in a matrix FCC phase. These phase differences are clearly evident on inspection of electron micrographs (magnification of 2,000 - 3,000 diameters) of the alloy. Different portions of the cold-worked sample, such as slices from a rod or bar, are then heated for about 8 hours each under temperature conditions differing incrementally by 50° or 100° F. over a range of temperatures between 800° F. and 1,500° F., depending on the alloy composition. Electron micrographs of each differently heat-treated portion are then examined. The temperature at which the HCP platelet phase begins to disappear from the original two-phased structure is the lower temperature limit of the transformation zone. The temperature at which all of the HCP platelet phase has disappeared, leaving only the FCC matrix phase, is the upper temperature limit of the transformation zone for an alloy of the particular composition studied. For the alloy of Example 1 below, for instance, the transformation zone lies approximately between 1,100° F. and 1,350° F.
Cold working of cast alloys at a temperature below the lower limit of the transformation zone may be by conventional practical means such as rolling, swaging, or drawing. Cold drawing or rolling are preferred, however, because they produce a more uniform product. When drawing to round shapes, the annealed materials are first submitted to centerless grinding to remove hot-rolling scale and to obtain a uniform bar size so that further area reduction by cold-working can be more readily controlled. After centerless grinding, the stock is suitably lubricated, for example with a coating of lime and/or molybdenum disulfide, and then cold-drawn.
Powdered alloys, prepared as described earlier herein may be handled in different ways. For instance, the powders may be hot- or cold-pressed into a desired shape and then sintered according to techniques known in powder metallurgy. The shaped articles can be cold worked by coining, for example. In another technique, the powdered alloy is sprayed hot (e.g. by so-called "plasma spraying") onto a substrate, such as of another metal, to which it adheres, and then cold worked in situ by suitable means such as swaging, rolling. or hammering.
Cold working is carried out to bring about a reduction in cross-section of at least about 5 percent. The most useful tensile strengths develop in alloys which have been cold worked to effect a reduction in cross-section of at least about 30 percent. As discussed earlier, the more the alloys are cold-worked, the more strength they develop. However, as the alloys increase in strength, they become less ductile. Accordingly, although the alloys can be worked to achieve a reduction in cross-section of as much as 95 percent, in practice they are not worked to an area reduction greater than 70 percent when the alloy is used, for example, for the manufacture of wire. Where still greater ductility is desired, as for example in alloys which are to be manufactured into fasteners, a reduction in cross-section between 50 - 60 percent is generally sufficient. In those alloys accordng to the invention which contain titanium with aluminum and/or columbium, desired high ultimate tensile strengths (260 - 280 ksi) can be obtained with a reduction in cross-section between about 35 and 45 percent. These alloys have particularly good ductility and retain it even at elevated temperatures.
To increase the strength of the cold worked alloys, the alloys may next be aged at a temperature between 800° F. and 1,350° F. for about 4 hours. A preferred aging temperature for those alloys free of titanium is about 900° F. If titanium is present, higher aging temperatures of about 1,200° F. are preferred. After aging, the materials are air-cooled.
It is pointed out that the alloys disclosed herein may also be worked at temperatures above the upper limit of the transformation zone (so-called "hot working") to produce articles of a desired shape. Such alloys, however, will not have the same high tensile strength characteristic of those alloys which are cold worked as hereinbefore described.
Also, it is possible to hot work those alloys of the present invention which have been priorly cold worked. Often, whatever loss of tensile strength is brought about by hot working is acceptable. Thus, for example, articles can be made by the partial hot forming (e.g. at about 1,900° F.) of previously cold worked stock. Although there is some loss in tensile strength in the hot worked portion, this may be acceptable if the hot worked portion is subjected to lower stresses than are those portions which have not been hot worked. The resultant article thus can combine desired strength properties with convenient and relatively low cost manufacture.
A better understanding of the present invention and of its many advantages will be had by referring to the following specific Examples, given by way of illustration.
EXAMPLE 1
An alloy, designated J80, having the following composition:
Percent by Weight C 0.01 Mo 9.9 Cr 19.4 Co 30.4 Fe 17.3 Ni 23.0 B 0.01
and having an electron vacany number, N v , of 2.52 was prepared from an approximately 50 pound charge comprising 9.90 lb of electrolytic chromium, 11.52 lb of nickel, 15.30 lb of cobalt, 5.10 lb of molybdenum, 0.04 lb of nickel-boron, and 8.70 lb of electrolytic iron. These components were melted in a MgO crucible in a vacuum furnace under 10 microns of mercury pressure at about 2,700° F. The alloy was refined until the leak-up rate, measured at 5-minute intervals, was substantially constant.
The melt was poured into a cylindrical steel mold with a copper stool, and then solidified. After cooling, the vacuum was broken, the furnace opened, and the mold stripped from the ingot, which was about 3.5 inches in diameter and 15 inches long.
Because the shape of the ingot permitted, the ingot was directly rolled at 2,050° F. to reduce its diameter to 0.75 inch, or an area reduction of about 22 to 1.
The rolled ingot was next annealed for 1 hour at 1,925° F., then centerless ground to a diameter of about 0.71 inch.
The ground piece was next drawn to 0.5 inch diameter using a lubricant of lime and molybdenum disulfide. The resulting 50 percent reduction in cross-section was achieved on about three passes, each effecting a 20 percent reduction in cross-section. The cold working was effected at ambient temperature, i.e. about 70° F.
The cold worked alloy was next aged for 4 hours at 950° F. and air-cooled. Specimens were then cut and ground for testing. The alloy showed the following properties. (ksi = kilopounds per square inch = 1,000 psi)
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 278.6 271.5 6.5 31.3 274.4 268.4 6.5 30.4
If the same alloy is cold worked to only a 38 percent reduction in cross section at 70° F., the following properties are found:
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 217.2 200.4 7.5 18.4 217.1 200.8 7.5 32.4
Although the alloys of the present invention contain significantly more iron than those taught in the aforementioned patent U. S. Pat. No. 3,356,542 to Smith, they surprisingly have a greater corrosion resistance than the prior art alloys.
In one test of corrosion resistance, a controlled voltage is applied by a potentiostat across a cell having a saturated calomel cathode and the alloy to be tested as the anode. A 3.5 percent NaCl solution is used as the electrolyte. The current through the cell is measured: by Coulomb's Law, the smaller the current observed, the smaller is the corrosion.
The J80 alloy of this Example, annealed and cold worked to an area reduction of 50 percent, showed the following corrosion characteristics in such a test:
Voltage Current (milliamperes) 0.2 nil 0.4 0.015 0.6 0.05 0.8 0.095
(Open circuit potential v. saturated calomel = 0.3 v.).
In contrast, a typical prior art alloy free of iron (whose composition is given below) showed the following results:
Voltage Current (milliamperes) 0.2 0.004 0.4 0.02 0.6 0.13 0.8 1.2
(Open circuit potential v. saturated calomel = 0.174 v.) This prior art alloy, which had been annealed and cold worked to an area reduction of 47.5 percent, had the following composition:
Percent by Weight C 0.02 Mo 10 Cr 20 Co 35 Ni 35
In another test for measuring the resistance of the alloys to crevice corrosion (i.e. corrosion under anaerobic conditions in which an oxygen concentration gradient is established, resulting in the formation of a concentration cell), bars of the J80 and prior art alloy were wrapped longitudinally with a stretched rubber band and immersed in 10 percent aqueous FeCl 3 solution for 72 hours at room temperature. No crevice corrosion was observed in any of the samples under these conditions. The degree of general surface corrosion was measured by determining weight loss per unit surface area.
Weight Loss (mg/in 2 ) J80 (cold worked 38 percent) 1.1 J80 (cold worked 50 percent) 1.2 Prior Art (cold worked 30 percent) 1.4 Prior Art (cold worked 30 percent) 3.2
(All four alloys were heat-aged for four hours at 1,000° F. after cold-working). Despite the high iron content of the alloy of the present invention, these tests show that both the crevice corrosion resistance and the general corrosion resistance of the alloy are, unexpectedly, comparable with the same properties of an iron-free prior art nickel-cobalt alloy.
EXAMPLES 2 - 7
A number of further alloy bodies were prepared as in Example 1 above having the composition shown below in Table I. While the percent by weight composition of each of the alloys falls within the ranges disclosed earlier herein, the N v value for certain of the alloys falls outside the permissible limits disclosed. Such alloys are not within the scope of the present invention and are shown for purposes of comparison. (All alloys contained 0.01 percent boron and 0.015 percent carbon in addition to the elements listed.)
TABLE I
Mo Cr Co Fe Ni N v J 83 10.0 20.0 31.5 9.5 29.0 2.40 J84 7.0 20.0 40.0 12.0 21.0 2.43 J88 8.5 20.2 30.3 19.3 21.7 2.54 J81 10.4 19.8 31.6 17.0 21.2 2.56 J82 10.8 20.3 32.8 16.7 19.4 2.61 J85 9.0 20.5 31.5 19.0 20.0 2.58
The alloys J81, J82, and J85 all contained a heavy sigma phase and were too brittle to be processed. Average properties for the remaining alloys after processing as in Example 1 above are reported in Table II.
TABLE II
Ultimate Tensile 0.2 Yield Reduction in Strength Strength Elongation Area Alloy (ksi) (ksi) (%) (%) J83 289.3 274.7 2.7 8.6 J84 287.7 256.9 5.2 21.9 J88 261.4 241.8 8.2 40.4
example 8
when preparing alloy bodies on a large scale, it is convenient to use ferroalloys in the melt. Thus, a 1,500 pound charge was prepared by vacuum melting the following components: 269 lb of low-carbon ferrochrome; 148 lb of pure molybdenum; 348 lb of nickel; 456 lb of cobalt; 88 lb of chromium; 190 lb of iron; 1 lb of nickel-boron; and 0.1 lb of additional carbon. A wet analysis of the alloy was as follows:
C 0.011 Cr 18.96 Si 0.01 Mo 0.86 Mn 0.03 Co 30.33 S 0.003 Fe 17.60 P 0.008 B 0.01 Ni Balance
EXAMPLE 9
The beneficial effects of heat aging after cold working are shown by comparing the properties earlier reported for the cold-drawn and heat-aged alloy of Example 1 with the properties (given below) for the same alloy after cold-working, but prior to heat-aging.
50 Percent Drawn (Average of Two Samples)
Ultimate Tensile 0.2 Yield Reduction Strength Strength Elongation in Area (ksi) (ksi) (%) (%) 218 205 8.8 41.0
38 Percent Drawn (Average of Two Samples)
195 150 13.5 45.5 EXAMPLE 10 Proceeding as in Example 1, the following alloy, designated as J153, was prepared by melting, casting, and rolling: Percent by Weight C 0.015 Mo 7.0 Cr 19.0 Co 35.7 Fe 9.0 Ni 25.5 B 0.015 Ti 3.0 Al 0.2 Cb 0.6
The effective atomic percent and N v for this alloy are calculated as follows.
The atom percent of each element present is first calculated, ignoring boron and carbon:
Atom Percent (No. of atoms per 100) Mo 4.25 Cr 21.29 Co 35.30 Fe 9.39 Ni 25.31 Ti 3.65 Al 0.43 Cb 0.38 Total 100.00
The number of hardener element ("X") atoms/100 alloy atoms is (3.65 + 0.43 + 0.38) = 4.46. The total unadjusted number of nickel and X atoms involved in Ni 3 X formation is (4 × 4.46) = 17.84, of which 13.38 are nickel, leaving (25.31 - 13.38) = 11.93 atoms of nickel available for matrix formation.
The number of molybdenum atoms which enter the Ni 3 X phase is the product of the Mo solubility in Ni 3 X, the atomic fraction of Mo in the alloy, and the number of Ni 3 X atoms, or (0.01 × 0.0425 × 17.84) = 0.01, leaving (4.25 - 0.01) = 4.24 atoms of Mo per 100 of alloy composition for matrix formation.
For chromium, (0.16 × 0.2129 × 17.84) = 0.61. Available Cr = (21.29 - 0.61) = 20.68 atoms/100.
For cobalt, (0.23 × 0.3530 × 17.84) = 1.44. Available Co = (35.30 - 1.44) = 33.86 atoms/100.
For iron, (0.15 × 0.0939 × 17.84) = 0.25. Available Fe = (9.39 - 0.25) = 9.14 atoms/100.
The total number of atoms, the number of atoms available for matrix formation, the effective atom percent, based on this latter total, and the effective atom fraction, are as follows:
No./100 Available Effective Effective (Atom %) No./100 Atom Percent atom Fraction Mo 4.25 4.24 5.31 .0531 Cr 21.29 20.68 25.90 .2590 Co 35.30 33.86 42.39 .4239 Fe 9.39 9.14 11.45 .1145 Ni 25.31 11.93 14.95 .1495 Ti 3.65 Al 0.43 Cb 0.38 Total 100.00 79.85 100.00 1.00
The N v for this alloy is 2.63, calculated by substituting these effective atom fractions into the equation
N v = 0.61 Ni + 1.71 Co + 2.66 Fe + 4.66 Cr + 5.66 Mo .
This value is less than the maximum permissible N v of 2.67 for an alloy containing 9 weight percent of iron from
N v = -0.017 W Fe + 2.82.
The alloy, thus, is substantially free of embrittling phases.
The alloy was next cold worked to reduce its cross-section by 40 percent, and then aged for 4 hours at 1,225° F. Samples of the alloy were tested at room temperature, 700° F., and at 900° F. to determine their strength and ductility, as reported in Table III below. Notched tensile strength (notch factor, K T = 6) and ultimate tensile strength were measured at each temperature and are compared in the Table.
TABLE III
Test Ultimate 0.2% Yield Temp. Tensile Strength Reduction in (°F) Strength (ksi) Elongation Area (ksi) (%) (%) R.T. 274.3 264.6 9.5 44.4 R.T. 268.7 256.6 11.0 43.8 Total 271.5 260.6 10.3 44.1 R.T. 379.4 NTS:UTS = 1.4:1 R.T. 380.4 Total 379.9 700 225.1 218.4 10.0 37.7 700 227.7 220.4 9.0 34.5 Total 226.4 219.4 9.5 36.1 700 320.8 NTS:UTS = 1.44:1 700 324.5 Total 322.7 900 205.6 201.2 10.0 33.0 900 202.8 197.6 11.0 37.3 Total 204.2 199.4 10.5 35.2 900 311.1 NTS:UTS = 1.50:1 900 299.8 Total 305.5
The effects of cold working on tensile strength and ductility can be seen by a comparison of Table III with following Table IV showing the same alloy properties in samples of the alloy of Table III cold worked to reduce their cross section by 46 percent.
TABLE IV
Ultimate Test Tensile 0.2% Yield Reduction Temp Strength Strength Elongation in Area (°F) (ksi) (ksi) (%) (%) R.T. 288.9 272.7 9.0 39.6 R.T. 282.5 273.6 8.5 36.2 Total 285.7 273.2 8.8 37.8 R.T. 388.5 NTS:UTS = 1.35:1 386.1 Total 387.3 700 226.5 220.4 8.0 26.7 700 232.3 222.2 9.0 33.3 Total 229.4 221.3 8.5 30.0 700 329.8 NTS:UTS = 1.46:1 700 332.8 Total 331.3 900 226.6 220.7 9.5 31.8 900 224.4 218.8 9.5 33.3 Total 225.5 219.8 9.5 32.6 900 317.7 NTS:UTS = 1.40:1 900 315.1 Total 316.4
EXAMPLE 11
Another alloy, designated J150, of lower iron content than the alloy of Example 10, was prepared as in Example 1. The composition of the alloy is given below. The N v of the alloy is 2.64, or below the permissible maximum of 2.70. The physical properties of samples cold worked to reduce their cross section by 40 percent and then aged at 1,225° F. for 4 hours are reported at various temperatures in Table V.
ALLOY COMPOSITION
Percent by Weight C 0.015 Mo 7.0 Cr 20.0 Co 36.2 Fe 7.0 Ni 26.0 B 0.015 Ti 3.0 Al 0.2 Cb 0.6
TABLE V
Test Ultimate 0.2% Yield Reduction Temp Tensile Strength Elongation in Area (°F) Strength (ksi) (%) (%) (ksi) R.T. 291.4 284.4 10.5 39.1 R.T. 286.3 282.3 10.5 42.9 Total 288.9 283.4 10.5 41.0 R.T. 382.9 NTS:UTS = 1.34:1 R.T. 392.6 Total 387.8 700 241.4 236.6 10.0 34.4 700 241.5 240.3 9.0 33.4 Total 241.4 238.5 9.5 33.9 700 350.2 NTS:UTS = 1.43:1 700 339.9 Total 345.0 900 240.0 230.5 4.0 8.6 900 237.7 228.5 5.0 10.0 Total 238.9 229.5 4.5 9.3 900 333.1 NTS:UTS = 1.38:1 900 329.3 Total 331.2
EXAMPLE 12
Following Table VI summarizes the nominal chemical composition of three alloys. Alloy A is a prior art iron-free alloy falling within the scope of U. S. Pat. No. 3,356,542 to Smith. Alloy B is an alloy according to the present invention, like the alloy of Example 1. Alloy C is an alloy like that in Example 10.
Table VII below compares the physical properties of these alloys at various temperatures after cold working to the same approximate degree of reduction in cross section and aging for 4 hours. ##SPC1## ##SPC2##