|Publication number||US3664889 A|
|Publication date||May 23, 1972|
|Filing date||May 26, 1969|
|Priority date||May 26, 1969|
|Also published as||DE2026957A1|
|Publication number||US 3664889 A, US 3664889A, US-A-3664889, US3664889 A, US3664889A|
|Inventors||Crooks Donald D, Fenn Raymond W Jr, Mccarthy William H|
|Original Assignee||Lockheed Aircraft Corp|
|Export Citation||BiBTeX, EndNote, RefMan|
|Patent Citations (7), Referenced by (11), Classifications (18)|
|External Links: USPTO, USPTO Assignment, Espacenet|
United States Patent McCarthy et a1.
 TERNARY, QUATERNARY AND MORE COMPLEX ALLOYS OF Be-AL  Inventors: William H. McCarthy, Los Altos; Raymond W. Fenn, Jr., Los Altos Hills;
Donald D. Crooks, San Jose, all of Calif.
Lockheed Aircraft Corporation, Burbank, Calif.
22 Filed: May26, 1969 21 Appl.No.: 827,752
 US. Cl. ..148/1l.5 R, 29/4205, 75/138, 75/150, 148/ll.5 A, 148/12.7, 148/32, 148/32.5  Int. Cl. ..C22f 1/04  Field of Search ..75/150, 138, 0.5 BB, 0.5 C; 148/1 1.5 F, 11.5 A, 11.5 R, 32,325, 12.7; 29/4205  References Cited UNITED STATES PATENTS 3,558,305 l/l97l Griffiths ..75/l5O 3,322,512 5/1967 Krock et a1. ..75/150 3,323,880 6/1967 Krock et a1. ..75/150 3,337,334 8/1967 Fenn et al. ..75/150 51 May 23, 1972 3,373,002 3/1968 Larsen et al. ..75/l50 3,379,513 4/1968 Larsen et a1. ....75/l50 3,438,751 4/1969 Krock et a1. ..75/l 50 Primary Examiner-Richard 0. Dean Attorney--Rodger N. Alleman and George C. Sullivan  ABSTRACT An alloy having extraordinarily high modulus to density and tensile strength to density ratios along with good workability and elongation properties is made from a beryllium-aluminum alloy powder having a complexity index between 1 and 5, as defined in U. S. Pat. No. 3,337,334. Magnesium or zinc or copper or one or more of several metals is blended with or deposited on the Be-Al powder; then the blend is consolidated without melting it and then worked, diffused, and heattreated, all below the eutectictemperature. The result is that the magnesium or other component (or components) of the alloy is dissolved substantially only into the aluminum-rich phase while leaving the complexity index unchanged. Hence, the properties are superior to alloys having the same Be-Al composition with or without other ingredients, especially those in which the magnesium or other metal is melted into the Be-Al composition.
24 Claims, 9 Drawing Figures GI-Be-36Al-3Mg looox Patented May 23, 1972 3,664,889
5 Sheets-Sheet 1 U 2 m N E Q 3 I LL 3 E 0' 5 T I E (D 3 LI... 3 5
INVENTORS WILLIAM H. MCCARTHY RAYMOND W. FENN, JR. DONALD D. ROOKS Patented STRENGTH, k si May 23, 1972 3,664,889
5 Sheets-Sheet 4 Be-35Al BASE ALLOYS TENSILE STRENGTH I ;2 YIELD STRENGTH Q so '5 (9 Z o .f a,
2o ELONGATION l I l l 1 2 4 6 8 IO l2 COMPOSITION OF ALUMINUM-RICH PHASE, WEIGHT PERCENT MAGNESIUM.
Be- 57AI BASE ALLOYS TENSILE STRENGTH i 2 6 '3 I g --Q (9 YIELD STRENGTH g 40 -20 ELONGATION l l l I l I o 2 4 e 8 l0 l2 I4 I6 0 COMPOSITION OF ALUMINUM-RICH PHASE, WEIGHT PERCENT MAGNESIUM.
INVENTORS FIG 8 WILLIAM H. MCCARTHY RAYMOND W. FENN, JR. DONALD D. CROOKS Agent Patented May 23, 1972 5 Sheets-Sheet 5 m QE llll lllmmlllllllllwllm 3 & x 25650 222 86 3% 23 m x m OOO Omo
INVENTORS WILLIAM H. MCCARTHY RAYMOND w. FENN, JR. DONALD D. CROOKS Agent TERNARY, QUATERNARY AND MORE COMPLEX ALLOYS OF Be-AL This invention relates to an improved alloy having three or more constituents, of which the principal constituents are beryllium and aluminum, and to a process for making the improved alloy.
The invention may be considered in general terms as an improvement on the alloy described and claimed in U. S. Pat. No. 3,337,334 issued Aug. 22, 1967. The terms defined in columns 21 and 22 in that patent are used in the same sense in this specification. According to that patent, binary berylliumaluminum alloys of superior mechanical characteristics can be prepared by atomizing a molten solution of aluminum in beryllium, the solution being initially at approximately 2,500 F. The rapid cooling inherent in the atomization process produces an intimate mixture of beryllium-rich and aluminum-rich phases in each atomized composite powder particle. When the process is properly carried out, the product has a complexity index of about 1 to 5 reciprocal microns, with aluminum-rich or beryllium-rich particle sizes of about i to 8 microns or less. When this binary alloy product is suitably consolidated and wrought, it has reasonably good tensile properties, as, for one example, a tensile yield strength of 45 ksi, a tensile ultimate strength of 57 ksi, and a 7 percent elongation, or, as another example with a somewhat higher weight percent of aluminum, a tensile yield strength of 33 ksi, a tensile ulti mate strength of 42 ksi, and an elongation of 7 percent. The product is extremely good when compared either with sintered mixtures of separate elemental powders of aluminum and beryllium in the same percentages or with composites that are prepared by casting, by liquid infiltration or by liquid phase sintering of aluminum or aluminum alloy into a beryllium powder mass, apparently because those other products do not have the fineness or the complexity of structure which are both essential to good mechanical properties.
in the present invention, additional elements are added to the binary beryllium-aluminum alloy of U. S. Pat. No. 3,337,334 without deteriorating that structure and in fact giving it even better properties.
It might be thought that complex beryllium-aluminum alloys could easily be obtained, simply by adding any desired additional elements to the Be-Al melt prior to atomization. However, this overlooks a critical fact: many of the most effective alloying elements have very high vapor pressure at the required melting temperature for the beryllium-aluminum base alloys, that is, at about 2,500 F. As a result, such addition elements evaporate before the melt can be atomized to form the powder. The following Table 1 presents some vapor pressure data for typical ternary addition elements and for the base constituents.
TABLEI Normal Boiling Point and Estimated Vapor Pressure at 2500 F. for Be, Al and Various Ternary Additions in Be-Al Alloys Approximate Vapor Pressure Normal Boiling Point at 2500" F.
Significant problems in the control of the composition of melts can be expected for all those elements having vapor pressure in excess of that for aluminum at 2,500 F, and this includes silver, bismuth, cadmium, lithium, magnesium, manganese, lead, and zinc.
in the process of achieving the present invention, we have 7 found it beneficial for the preponderance of certain addition elements to be partitioned in the solid alloy to the aluminumrich phase rather than to the beryllium-rich phase. Melting of ternary and higher order compositions tends to distribute the addition elements more or less equally between these two phases. Inferior properties may result even with those addition elements that do have low vapor pressure, such as chromium, copper, iron, and silicon, because of the undesirable degree of partition between the two phases.
In the present invention, binary beryllium-aluminum alloys in powder form having satisfactory microstructure are first prepared, as by the melting and atomization process described in U. S. Pat. No. 3,337,334. The product is not remelted subsequently; in this invention remelting is specifically avoided so that the fine complex microstructure is preserved, and so are the excellent mechanical properties, which can be substantially improved by subsequent alloying according to the instant process. Instead of melting the beryllium-aluminum alloy, the binary alloy powder is blended with powdered addition elements, which are mixed thoroughly to achieve good distribution. High vapor pressure elements, such as those mentioned before in connection with Table l, and especially magnesium, have been found to be very beneficial addition elements. The powder blend is consolidated by a suitable hot or cold operation, carried on without any melting. For example, hot pressing or explosive compaction may be used. The resultant blend may then be worked, as by extrusion or sheet rolling, to produce a semi-finished article and given a diffusion annealing treatment appropriate to the addition elements which are present in order to distribute these elements evenly in solid solution throughout the aluminum-rich phase in the structure. Some or all of the diffusion annealing treatment may precede the working, and none or all of it may follow the working, depending upon the composition being prepared. Temperatures may be selected to control the rates of diffusion of the addition elements into the aluminum-rich phase relative to those in the beryllium-rich phase, thereby achieving a desired partition between the phases. For example, we have found that magnesium partitions substantially to the aluminum-rich phase only and that this partitioning is beneficial, in that the magnesium strengthens and increases the workhardening coefficients of this softer phase and results in strengthening the composite when compared to a binary beryllium-aluminum alloy.
The compositional diffusion of the aluminum-rich phase containing addition elements can be homogenized by hot or cold working. The strained structure can also accelerate diffusion. The complexity of the microstructure of the binary beryllium-aluminum composite powder or the consolidated article is very beneficial in promoting diffusion of the alloying element or elements, because of the very large area of interphase and grain boundary surfaces per unit volume, which accelerate the diffusion, especially at lower temperatures. It is important that the diffusion heat treatments be performed at temperatures below the lowest eutectic temperature in the alloy system considered; by doing this, eutectic melting is avoided. Eutectic melting tends to increase phase particle size and grain size, to introduce voids, to segregate the ternary addition, and to decrease the complexity of structure, and, as a result, it reduces the mechanical properties. For example, properties little or no better than those which are obtained by liquid phase processing can result if there is eutectic melting.
The present invention may be compared with some of what has been done in the prior art. For example, liquid infiltration, as described in U. 5. Pat. No. 3,323,880, produces a coarse grain structure of insufiicient complexity and having inferior mechanical properties, especially in thicknesses exceeding one inch. The alloy shown in FIG. 3 of U. S. Pat. No. 3,323,880, for example, has a mean free path D in the beryllium-rich phase of 27 ,u and A, in the aluminum-rich phase of 1 l u, with a complexity index of 0.28 t. These values may be compared with a D of4 p., an A, of 2 ,u., and a complexity index of 2 p. of the alloys of this invention. Such an alloy also suffers structurally from the fact that the aluminum-rich phase contains some beryllium-aluminum eutectic material, which is more or less dispersed in the aluminum-rich phase, which also contains some magnesium or other addition element. On the other hand, the splat cooling or atomization technique of U. S. Pat. No. 3,337,334 can be used to provide a suitable binary alloy (Be-Al) for use as a raw material in the present invention. The present invention is to be distinguished from alloys (See U. S. Pat. No. 2,994,947) where aluminum-containing intermetallic compound formation is desired and solid-solution formation deliberately avoided. The present invention involves solid-state diffusion of the addition metal or metals into Be-Al master alloys. As a result, the large amount of interphase boundaries in the binary alloy produced according to U. S. Pat. No. 3,337,334, greatly enhances the mass transport, especially at lower temperatures.
An example of what can be done in the present invention is given in Table II, where there is a comparison between certain properties obtained by the present invention, by the aluminum-beryllium alloy of U. S. Pat. No. 3,337,334, and by liquid phase processing.
TABLE II Typical Mechanical Properties of Beryllium-Aluminum Base Alloys Compo- Tensile Tensile sition Yield Ultimate Method of wt. Strength Strength Elon- Preparation Al KS1 KS1 gation l. Consolidation of complex particles per U.S. Pat. No.
3,337,334 (Be Al only) 35 45 57 7 2. Ternary (BeAl-Mg) alloy of this invention prepared from Item I 34 72 93 7 3. Consolidation of complex particles per U.S. Pat. No.
3,337,334 (Be-Al only) 57 33 42 7 4. Ternary (Be-Al-Mg) alloy of this invention prepared from Item 3 54 56 74 6.5 5. Consolidated mixtures of separate elemental powders (Be-Al only) 41 39 44 0.6 6. Liquid phase sintering of Al alloy into Be powder 30 25 40 1 7. Fine-grained cast Be-ALMg alloy 30 28 35 4 Other objects and advantages of the invention will appear from the following description of some preferred embodiments thereof.
In the drawings:
FIG. 1 is a photomicrograph (100 magnification) of a typical 6lBe-36Al3Mg alloy extrusion embodying the principles of this invention.
FIG. 2 is a photomicrograph (1,000 mag.) of the 61Be- 36Al3Mg alloy shown in FIG. 1.
FIG. 3 is a photomicrograph (100 mag.) of a 63Be-37Al binary alloy extrusion prepared from the same Be-Al powder that was used to make the ternary alloy of FIG. 1.
FIG. 4 is a photomicrograph 1,000 mag.) of the 63Be-3 7A1 alloy extrusion shown in FIG. 3.
FIG. 5 is a graphical diagram showing the quotient of tensile strength divided by density for some of the alloys of this invention and some other structural materials.
FIG. 6 is a graphical diagram showing the quotient of the square root of Youngs modulus divided by the density of some of the alloys of this invention compared with some other structural materials. This quotient is a figure of merit for materials to be elastically loaded in compression, and its reciprocal is known as the elastic weight index.
FIG. 7 is a graph of the tensile properties of the Be-35Al-Mg base alloys of this invention versus weight percent magnesium in the aluminum-rich phase.
FIG. 8 is a similar graph for the Be-57Al-Mg base alloys of this invention.
FIG. 9 is a phase diagram for alloys near the Be-Al edge of the Be-Al-Mg system.
EXAMPLE 1 Starting Materials Different lots of binary beryllium-aluminum alloy powder were used as the base material in Example 1. These were 40 mesh Be-Al and Al-Be alloy powders supplied by the Beryllium Corporation and manufactured by the Lockalloy process (U. S. Pat. No. 3,337,334). This powder preparation process involves the inert gas atomization of a stream of molten beryllium-aluminum alloy heated originally to about 2,500 F. The very rapid cooling inherent in this method produces reasonably fine particles containing both beryllium-rich and aluminum-rich phases, intimately mixed in each powder particle as a divorced eutectic microstructure. The mean free path in the aluminum-rich phase is about 2 microns, and in the beryllium-rich phase, about 4 microns. One type of powder contained nominally 35 to 37 percent aluminum, which is within the standard composition range for commercial Lockalloy, and the second type had a nominal aluminum content of 57 percent. The powders contained about 0.2 percent each of BeO and A1 0 500 ppm carbon, 600 ppm iron, 150 ppm nickel and 500 ppm silicon. Other metallic impurities were less than ppm.
The materials added to the Be-Al binary alloy powders were usually elemental metallic powders. The powders were added with the object of obtaining solid solution and/or precipitation hardening in the aluminum-rich phase of the composite. Magnesium was a Baker and Adamson Code 1900 commercial grade powder of about -40 mesh particle size. Zinc was 99.999 percent -325 mesh powder produced by Cominco Products, Inc. The copper powder was commercial grade Baker and Adamson Code 1618 material of about -200 mesh particle size. Lithium was added as hydride of about 99 percent purity and 200 mesh grain size which was obtained from F oote Mineral Company.
Alloy Preparation The composites in Example 1 were prepared from binary beryllium-aluminum alloy powders containing either 35 to 37 percent or 57 percent aluminum. Ternary and higher order alloying elements were added to the base alloy powder. The mixes were vacuum hot-pressed, and some diffusion annealing operations were carried out in the vacuum hot-press. The resulting compacts were extruded into round bars. Additional diffusion and strengthening heat treatments were applied. Tensile bars were prepared from the extruded material. The merit of the various alloying elements was assessed on the basis of room temperature tensile strength, ductility and elastic modulus as compared with similar properties of the binary compositions which were prepared for control purposes. Details of the preparation of some of the various alloys are given in Table IV.
Powder Blending The ternary and quaternary additions were blended with the beryllium-aluminum powder in a blender like a household blender. A typical batch contained 60-120 grams of Be-Al powder, together with the appropriate weight of addition material. This was shaken by hand in the one pint mixing bottle to distribute the addition material, and then it was mechanically blended for approximately 30 seconds. A further addition of about 0.3 ml. per gram of solids of isopropyl alcohol was often made, and the material was again blended for approximately one additional minute. The purpose of the isopropyl alcohol was to desorb oxidizing gases from the surfaces of the powder particles and thus promote sintering in operations to follow and possibly to minimize oxide impurity content. The product was a thoroughly mixed moist powder of about the consistancy of foundry molding sand, when the alcohol addition was used. The powder mixture was packed by hand into a copper cylinder which was sealed at one end. A copper plug was then inserted in the can following the powder charge. The copper can was a convenient container for transferal to the hot press, but it served principally as an enclosure and lubricant during the subsequent extrusion operation.
Hot Pressing The l or 1% inch nominal outside diameter copper pipe containing the powder charge was placed in a close fitting, 2 outside diameter steel sleeve to provide radial support during hot pressing. The assembly was inserted in a vacuum furnace provided with a 2 7/16 inch inside diameter inconel tube. The hot pressing load was applied to a punch inserted in the copper pipe on top of the copper plug. Control of temperature was achieved by the use of thermocouples and temperature controllers. A vacuum system, including a mechanical forepump, 2-inch diffusion pump and liquid nitrogen cold trap was used. The system was evacuated with forepump only, until the isopropyl alcohol had evaporated, and then the diffusion pump was applied to the system. When the pressure was below about 8X10 mm. mercury the furnace was heated to 800 F. For alloys containing magnesium, zinc or lithium, the diffusion pump was cooled and the pressure allowed to rise to about 30-50 microns, before the temperature reached the point that significant loss of the more volatile addition elements could be expected (e.g., approximately 500 F for magnesium). In the case of lithium, lithium hydride was added, and a thermal arrest accompanied by a sharp increase in pressure was observed at about 650 F upon heating. This was the point where the hydride decomposed and hydrogen was produced.
Upon reaching 800 F, a load of about 15,000 to 25,000 pounds was applied and held for approximately 30 minutes. This produced a nominal stress of about 10,000 to 25,000 psi in the powder mass and densities in excess of 90 percent were achieved. The deflection of the crosshead was measured during the pressing operations and it was found that substantially all of the reduction in volume occurred within minutes after reaching 20,000 pounds. The pressed billet was allowed to remain at temperature and under vacuum for several hours to begin compositional homogenization. Stress was again applied to the pressing for about 15 minutes, after which the pressing was allowed to cool in vacuum to some temperature below 300 F. During the hot pressing operation, the copper plug on top of the powder charge would deform around the punch. This, together with a small amount of powder which extruded out between the plug and copper can, made a very effective seal for the top of the can. Thus, when the copper can was pressed out of the steel supporting sleeve, a vacuum encapsulated extrusion billet was obtained. It was only necessary to trim off the excess copper and point the end of the pressing. Re-encapsulation of the hot pressing, prior to extrusion, was completely avoided.
Extrusion The 1.25 or 1.50 inch diameter copper clad extrusion billets were extruded at temperatures from 600 F to 950 F through a 0.385 or 0.491 inch diameter die, the temperature depending upon the composition. The extrusion ratio was thus nominally 10:1. The breakthrough force required varied from 55 to 130 tons. Running load was about 90 percent of the breakthrough force, to give extrusion constants of 35,000 to 65,000 psi. (The extrusion constant is where P is the running extrusion force, A0 is the inside area of the container and R is the reduction ratio.) The billets were heated for about 30 minutes in an air furnace or a salt bath, transferred to the extrusion tools (which were preheated to 400 to 600 F) within 10 seconds and extruded within about 2 minutes of removal from the furnace. It was found that extrusion speed was an important factor to consider and that ram speeds greater than about 15 inches per minute produced inferior material. Table 11] below gives some tensile properties versus extrusion speed for round bars. The extrusion direction was opposite to the hot pressing direction. Each billet produced about 18-30 inches of extruded bar. Some 0.28 by 1.08 inch round-edge flat bars were also produced. The details of consolidation and extrusion are given in Table IV, later.
TABLE III Tensile Properties of Be-36Al-3Mg Alloy as a Function of Extrusion Ram Speed. (10:1 extrusion ratio, through 0.491 inch diameter die.)
Sample Preparation and Heat Treatment The copper was removed from the extruded bars by lathe turning and/or nitric acid etching. Three-inch-long coupons for tensile specimens were cut from the extrusions. The coupons, 0.30 to 0.45 inch in diameter, were diffusion annealed in air for times and temperatures appropriate to the particular composition used. The coupons were quenched in water directly from the diffusion annealing furnace. Some samples were given subsequent aging treatments, e.g. hours at 800 F was a typical treatment for Be-Al-Mg alloys.
Evaluation The principal basis for evaluation of the materials was the room temperature tensile test. Tensile specimens with conical ends for gripping were machined after heat treatment. Reduced sections were usually 0.14 to 0.25 inch in diameter by 1.5 inches long, but occasionally the limited availability of material necessitated smaller reduced sections. Some threaded-end tensile specimens were also used. Approximately 0.002 inch per surface was removed by hot NaOl-l etching from some of the specimens after machining. This surface condition is necessary to achieve good tensile ductility in commercial beryllium, but was found to be unnecessary for Be-Al and for Be-Al base alloys. Youngs Modulus determinations were performed using dead weight loading and a pair of Tuckerman optical extensometers. The sensitivity of this strain measuring system is about 2 microstrain. The specimens could be loaded elastically to at least 5,000 psi so that the strain measurement sensitivity was better than 1 percent of the maximum elastic strain. To measure strength and ductility, the specimens were loaded in a tensile testing machine at a crosshead travel rate of 0.01-inch per minute. This resulted in strain rates of about 6-8 10" per minute during the elastic portion of the stress-strain curves and about 3-5X10' per minute during plastic deformation. Strain was measured with an averaging pair of Baldwin Model T2M Microformer extensometers. Tensile strength, 0.2 percent offset yield strength,
36-Al-3Mg alloy was found to have a mean free path D of 3.9 ,u. in the beryllium-rich or gray phase, a A, of 1.6 p. in the aluminum-rich or white phase, and a Complexity Index of 2.0 a. The 63Be-37Al binary alloy extrusion of FIG. 4 has a mean free path D of 3.9 p. in the beryllium-rich phase (gray), a A, of 1.6 p. in the aluminum-rich phase (white) and a Complexity Index of 1.9 t. Details of Example 1 Processing details for examples of 16 different alloys, two of them Be-Al, l0 Be-Al-Mg alloys, and one each of Be-Al-Cu, Be-Al-Li, Be-Al-Zn, and Be-Al-Mg-Zn, were prepared according to the above described process and that shown in Table IV below. Tensile test results obtained on 12 different Be-Al and Be-Al-Mg alloys are shown on Table V, which also includes known values of various other alloys and metals, which are given for comparison, including some commercial Be-Al alloys, Be itself, Mg and Mg alloys, Al, Ti-Al-V alloys, steels and stainless steel. The results are tabulated below, and discussion follows:
Discussion The alloys of Table IV containing magnesium had microstructures, after diffusion annealing for 100 hours at 800 F, :followed by water quenching, which were substantially identical to those of the binary Be-Al compositions. The locations of the 'Mg particles in the original powder mix were converted to an aluminum-rich solid solution. Some of these beryllium-free regions are marked 2 in FIG. 1. In our example, these regions were less than 0.005 inch (l00 p.) thick. Magnesium has diffused away from these regions, and aluminum has diffused into them, leaving a continuous structure. No intermetallic compounds were found by X-ray diffraction; and lattice parameter measurements on the aluminum-rich phase indicate a Mg-Al solid solution containing the nominal starting concentration of magnesium. FIGS. 1 and 2 show the micros- .tructures of a typical Be-36Al3Mg alloy. These can be compared with FIGS. 3 and 4, which show the respective microstructure of a binary Be-37Al extrusion that was prepared for control purposes.
Results of the room temperature tensile test on some of the Be-Al-Mg alloys and other materials are given in Table V. These data show that very substantial increases in strength can be achieved by magnesium additions, without serious loss in TABLE I V.ALLOY PREPARATION Hot pressing Extrusion Max. gas Total pressure, Max. time at Max Extrusion Extrusion Composition (remainder microns, temp., temp., stress, 'Iernp., constant idcnitfication beryllium) wt. percent Hg F. hrs p.s.1 F. p.s.i.
a A1 0.4 1,120 7. 3 34,900 950 51,500 b. 35 A11.8 Mg 20. 850 19. 4 23, 800 800 51, 500 c 36 A1-2. 7 Mg 40. 810 20. 3 10, 400 650 64, 500 (1 34 Al3.4 Mg 12. 820 70. l 23, 600 800 41, 300 0 34 A13.7 Mg 50. 810 15.5 23, 600 800 51, 500 f 33 A1-5.4 Mg 15. 810 16. 1 24,200 800 46, 800 57Al 0.4 840 1.1 ,600 800 36,500 11 57 A11.2 Mg 35. 860 19. 1 24, 000 800 43, 400 i 56 Al3.0 Mg 20. 810 15. 4 23, 400 800 58,500 J A14.1 Mg 25. 820 20.0 17, 900 800 47, 500 k 54 A16.0 Mg 8. 810 0. 3 23, 400 800 44, 700 53 A1-8.6 Mg 20. 760 20.9 23, 200 800 44,700 56 Al2.7 Cu 0. 1 950 118. 5 23, 900 800 47, 800 56 A1-1.0 L1. 2,000. 800 18.5 24, 300 800 34,700 34 111-1. 7 Zn 35. 715 21. 6 23,800 800 47,500 p 33 Al1.7 Zn1.7 Mg 60. 715 21.1 24, 200 800 48,100
TABLE \'.TENSILE PROPERTIES AND STRUGTURAL EFFICIENCY PARAMETERS FOR Be-Al-Mg ALLOYS COMPARED WITH OTHER AEROSPACE MATERIALS Composition e Representative tensile properties b Structural efliciency Percent Al For stiflness in binary Yield Tensile Elongation, Young's Density a For strength, Be-Al Percent Mg Percent Mg strength, strength, percent in modulus (E) (theoretical) /P 10 TS/ 10 Extrusion powder in alloy in Al phase s.i. (TS) K s.i. 1 inch 10 p.s.i. p lb./in. lb.- in. inches 35 0 0 44. 61. 12. 30. 4 0. 0749 74. 0. 81 35 1. 8 5. 3 74. 93. 10. 30. 4 0.0749 74. 1. 24 36 2. 7 7. 1 77. 97. 8% 30. 3 0 0748 74. 1. 30 35 3. 4 10. 1 73. 93. 7. 30. 2 0. 0747 74. 1. 24 35 3. 7 11. 1 70. 72. 1% 28.6 0. 0747 74. 0. 96 35 5.4 16.3 76. 82. 2. 28.6 0.0747 74 1. 09 57 0 0 34. 42. 8 23. 0 0. 0814 58 0. 51 57 1. 2 2. 0 48. 56. 10. 22. 1 0.0814 58 0. 69 57 3. 0 5. 3 47. 68. 15. 22.3 0. 0812 58 0. 84 57 4. 1 6. 9 56. 79. 14. 22. 3 0. 0810 58 0. 98 57 6. 0 11. 1 56. 74. 6% 22.4 0 0807 59 0. 91 57 8.6 16. 4 56. 73. 7 20. 6 0. 0798 59 0. 91 62-Be-38Al; Lockalloy extrusion, annealed 44. 56. 5 28. 0. 0756 70. 0. 62-Be-38Al; Lockalloy sheet, annealed... 37. 50. 5 28. 0.0756 70 0. 67 Be; Hot pressed block 41. 49, 1 42. 5 0. 067 97 0.73 Be; powder sheet 54. 75. 5 42. 5 0. 067 97 1. 12 Mg; HK 31A-H24 Sheet 29. 36. 4 6.5 0. 0647 39 0.56 Mg; ZKGOA-T5 extrusion- 40. 52. 5 6. 5 0. 0659 39 0. 78 Mg-Li; LA141-T5 Sheet 20. 24. 18 6. 2 0.0485 51 0. 50 Al; 7075-T6-c1ad sheet 63. 72. 7 10. 3 0. 101 32 0. 72 Ti; 6Al, iv-annealed- 137. 147. 10 16. 2 0. 25 0. 92 Ti; 6A1, 4Vso1ution treated and aged 161. 178. 5 16.2 0.160 25 1. 11 Steel; 300M-forgings 240. 300. 8 30. 0. 283 19 1. 06 Steel; 18 Ni maraging (200)-bar 209. 214. 12 26. 2 0. 289 18. 0.74 Stainless steel, PH15-7 Mo cond. RH950-sheet 225. 240. 6 28. 0. 277 19. 0.87
I PlgrcentlMlgl in Al-rich phase according to the made-up composition and assuming that all of the Mg partitions to the Al-rich phase, leaving none in e e-ric p ase.
b Tensile properties are averages of two to four tests, except that single values are given occasionally when the second test result seemed abnormal due to some random defect. Values for materials other than Be-Al base alloys are typical of commercial production, except that some ductility values are specification minimums.
The stifiness structural efficiency is the reciprocal 0! Elastic Weight Index and is a useful figure oi merit for many compression applications. The sgren7gil11 ls ltrucltural efi-lclency value is a useful figure of merit for tensile application, E=30.4Xl0 p.s.i. for Be-35Al-Mg alloys and E=22.4X10 p.s.1. for
e-5 g a 0y.
4 Density is on the assumption that all the magnesium content is in the aluminum-rich phase.
corresponding tensile yield strengths.
TABLE VI Compression Properties of Be-Al-Mg Alloys at Room Temperature (Heat Treatment: 800 F., 100 hours, Water Quenched) Table V and FIGS. 5 and 6 demonstrate the unique combinations of properties which can be obtained by this new process and, by way of example, utilized in the Be-Al-Mg alloys. The percent Al base alloys containing Mg have strength to density ratios of up to 1.30 X 10 inches which surpasses all other aerospace structural materials by at least 16 percent. These Be-Al-Mg alloys are, in fact, 81 percent better, on a strength/density basis, than the widely used aluminum base alloys. Simultaneously, the excellent stiffness structural efiiciency of commercial Lockallo'y is maintained at levels 40 percent greater than all other materials, except pure beryllium. The Be-35Al-3Mg composition is 180 to 250 percent more efficient, on a stiffness basis than the strong steel and titanium alloys which approach the strength-to-density ratio of the Be-Al-Mg alloy.
The 57 Al base alloys, while significantly inferior to pure Be and the Be 35 percent A] base alloys, are structurally more efficient than all other materials, and in addition, should have substantially better ductility and formability than the other berylliferous alloys. That is, Be-54Al-6Mg has stiffness structural efficiency better than magnesium-lithium alloy (which, in turn, is substantially better than any other alloy not containing Be). The Be-55Al-4Mg alloy has a strength-to-density ratio comparable to those of the strongest beryllium, steel and titanium alloys. Its strength-to-density ratio is 26 percent better than the best magnesium base alloys and 36 percent better than 7075-T6 aluminum alloy.
Mechanical properties are plotted versus magnesium content in FIGS. 7 and 8. These figures show that strength leveled off after magnesium content in the aluminum-rich phase exceeded about 8 percent. In these tests, further magnesium additions did not greatly improve strength; but they did degrade ductility. This situation may not be an inherent characteristic of the Be-Al-Mg system, but rather, it may be due to inadequate homogenization of the Mg content. Thus, there may be regions, near the original locations of Mg powder particles, which are deficient in Be (See FIG. 1). Preliminary results of an assessment of magnesium concentration by lattice parameter measurements using X-ray diffraction suggest that this may be so. Thus, it may be possible to substantially improve 6 strength, and to increase ductility, by homogenization at higher temperatures, or for longer times or by use of finer Mg additions to shorten the diffusion path length, so that the Mg is even more uniformly distributed in the Al-rich phase. The initial slope of the strength versus magnesium content curves in FIGS. 7 and 8 may, under these conditions, extend out to about 10 percent or more Mg, with superior results.
EXAMPLE 2 An alloy composite containing nominally 36 percent Al, 3
was prepared by vacuum vapor deposition of magnesium coatings onto sheets of a binary beryllium-aluminum alloy consisting of a fine, complex mixture of beryllium-rich and aluminum-rich phases containing initially 38 percent Al, 62 percent Be. Sheets of suitable binary beryllium-aluminum alloy produced according to U. S. Pat. No. 3,337,334 are available commercially under the trade name Lockalloy. The coated sheets were vacuum annealed at 800 F for suitable time to achieve a uniform alloying of the aluminum-rich phase with magnesium, to produce an aluminum-magnesium solid solution, without disturbing the fineness or complexity of the beryllium-rich and aluminum-rich phases in the starting sheets of Be-Al alloy. The Be-Al-Mg alloy composite so obtained had the following mechanical properties: 41,000 psi yield strength, 58,000 psi ultimate tensile strength, 6 percent elongation. The binary Be-Al starting sheets had the following mechanical properties which are typical of material prepared according to U. S. Pat. No. 3,337,334; 36,000 psi yield strength, 48,000 psi ultimate tensile strength, 8 percent elongation.
EXAMPLE 3 In addition to magnesium, other addition elements are also possible in this invention. In each case, they form part of the aluminum-rich phase. Theoretically, such addition elements may be dissolved in amounts up to the maximum amounts shown in TABLE VII, as follows:
TABLE VII Maximum Limits of Solid Solubility for Various Binary Additions to Alpha Aluminum Alloys.
Titanium, chromium, bismuth, cadmium, lead, zirconium, iron, and nickel are elements of moderately low solubility in aluminum. Accordingly, they are of little value as ternary additions to the Be-Al base alloys used in this invention. However, they are of some importance as quaternary additions, or
as additions in more complex alloys built on Be-Al base alloys, especially when these contain zinc, magnesium, copper, lithiurn, silver, silicon or manganese, or combinations of these just-named elements.
An alloy composite containing 37 percent Al, about 4 percent Zn, balance Be with traces of BeO, A1 0 ZnO was prepared by vacuum vapor deposition of zinc coatings onto sheets of a binary beryllium-aluminum alloy consisting of a fine, complex mixture of beryllium-rich and aluminum-rich phases containing initially 38 percent Al, 62 percent Be. Sheets of suitable binary beryllium-aluminum alloy, produced according to U. S. Pat. No. 3,337,334, are available commercially under the trade name Lockalloy. The coated sheets were vacuum annealed at 670 F to l,070 F for sufficient time to convert the aluminum-rich phase into an aluminum-zinc solid solution, without disturbing the fineness or complexity of the beryllium-rich and aluminum-rich phases of the starting sheets of Be-Al alloy. To avoid eutectic melting and consequent deterioration of this structure, it is necessary to limitof a beryllium-rich phase and an aluminum-zinc solid solution phase, the latter containing up to 14 percent zinc.
EXAMPLE 4 Beryllium-aluminum based alloys containing 5-80 percent Al, 0.1-1.0 percent Mn, balance beryllium are prepared by mixing binary Be-Al powder with a suitable amount of Mn powder, the amount of Mn being selected so that it is about 0.3 to 1.85 percent of the aluminum plus manganese content. A typical alloy could contain 40 percent A], 0.5 percent Mn, balance Be impurities. (The Be-Al powder can be made according to the teaching of U. S. Pat. No. 3,337,334 and the processing described in this example does not significantly alter the fineness nor complexity of such Be-Al powder, but does convert the aluminum-rich constitutent to an Al-Mn solid solution.) After consolidation of the Be/Al Mn mixed powder, and suitable solid state diffusion heat treatments together with mechanical working, the three phase mixture is converted into a two phase alloy. This alloy has a microstructure consisting of a fine, intimate mixture of Be-rich particles, juxtaposed with Al-Mn solid solution particles. The strength of the ternary composite, after solution heat treatment, is about 5-20 percent greater than that of the corresponding binary Be-Al alloy at room temperature. At 500 the strength improvement is about -50 percent, depending upon the beryllium content.
EXAMPLE 5 The alloys described in Example 4 can be prepared in an alternate manner by atomizing a melt of Be+Al+Mn at approximately 2,500 F according to the teaching of U. S. Pat. No. 3,337,334. Results arequite similar.
EXAMPLE 6 EXAMPLE 7 The Be-Al-Si alloys described in Example 16 can be prepared in an alternate manner by atomizing a melt of Be Al Si at approximately 2,500 F according to the teaching of U. S. Pat. No. 3,337,334. Results are similar to those of Example 6.
EXAMPLE 8 Beryllium-aluminum base alloys containing 5-80 percent Al, 0.02-0.80 percent Si, 0.03-l.2 percent Mg are prepared in a manner analogous to that for the Be-Al-Mn alloy described in Example 4. The amounts of Si and Mg are selected so that (Si Mg is approximately 0.3 to 1.85 percent of the total of (Al Si Mg). It may be preferable that the Si content be slightly in excess of that required to form the stoichiometric compound M g Si. A typical alloy could contain 40 percent A1, 0.3 percent Mg, 0.2 percent Si, balance Be impurities. The quaternary alloy, in the solution heat treated condition, has
room temperature strength about 5-20 percent greater than that of the corresponding binary Be-Al alloy. The Be-Al-Mg-Si alloy can be further strengthened by precipitation hardening to 10-30 percent improvement over the Be-Al. In either solution treated or precipitation hardened condition, the Be-Al-Si- Mg quaternary alloy is stronger than the corresponding Be-Al alloy at 300 F by approximately 15-40 percent, depending upon beryllium content.
EXAMPLE 9 The Be-Al-Si-Mg alloys described in Example 8 can be prepared by first atomizing an Al+Be+Si melt at about 2,500 F according to the teaching of U. S. Pat. No. 3,337,334. To this ternary Be-Al-Si powder is then added Mg powder and the mixture is consolidated and prepared by the methods described in Examples 4 and 8. (Silicon and manganese can be added to the melt at atmospheric pressure, but magnesium cannot.) Results are similar to those obtained for Example 8.
EXAMPLE 10 Beryllium-aluminum based alloys containing 5-80 percent Al, 0.5-10.0 percent Mg, 0.02-l.0 percent Mn, 0.005-0.3 percent Cr are prepared in a manner analogous to that for the Be-Al-Mn alloy described in Example 4. The amount of Mg is selected so that it is about 0.5-l3 percent of the total of (Al+Mg+Mn+Cr). Mn is approximately 0.5 to 1 percent and Cr is about 0.01 to 0.3 percent of the (A1+Mg+Mn+Cr) total. A typical alloy could contain 40 percent Al, 2 percent Mg, 0.05 percent each of Mn and Cr, balance Be impurities. The quintary Be-Al-Mg-Mn-Cr alloy is about 10-30 percent stronger than its binary Be-Al counterpart at room temperature, and 20-100 percent stronger at 500 F, depending upon Be content.
EXAMPLE 1 l The Be-Al-Mg-Mn-Cr alloys described in Example 10 can be prepared in an alternate manner by first atomizing a melt of Be+ Al+ Mn+ Cr at approximately 2,500 F according to the teaching of US. Pat. No. 3,337,334. The Mg is then added as powder and the process carried out as in Example 10, with similar results.
EXAMPLE 12 Beryllium-aluminum based alloys containing 5-80 percent Al, 0.5-6.0 percent Cu, 0.1-4.0 percent Mg, 00-10 percent Mn are prepared in a manner analogous to that for the Be-Al- Mn alloy described in Example 4. The amounts of Cu, Mg and Mn are selected so that they are about 1.0-6.0 percent, 00-10 percent and 0.5-4.0 percent, respectively, of the total (Al+Cu+Mg+Mn) content. A typical alloy could contain 40 percent A], 2 percent Cu, 0.6 percent Mg, 0.2 percent Mn, the balance Be and impurities. This Be-Al-Cu-Mg-Mn alloy, in the solution heat treated condition, is about 20-200 percent stronger than its binary Be-Al counterpart; 30-200 percent stronger in the precipitation hardened condition. At 300-5 00 F, the Be-Al-Cu-Mg-Mn alloy is about 30-300 percent stronger than the corresponding Be-Al alloy, depending upon Be content.
EXAMPLE 1 3 The Be-Al-Cu-Mg-Mn alloys described in Example 12 can be prepared in an alternate manner by first atomizing a melt containing some of all of the Be+Al+Cu+Mn at approximately 2,500 F according to the teaching of U. 8. Pat. No. 3,337,334. The Mg and the remaining components are then added as powder and the process carried out as in Example 12, with similar results.
EXAMPLE l4 Beryllium-aluminum base alloys containing 5-80 percent Al, 0.03-7.0 percent Mg, 0.02-1.5 percent Si, 0.01-3.0 percent Cu, 0.000-0.3 percent Cr are prepared in a manner analogous to that for the Be-Al-Mn alloy described in Example 4. The amounts of Mg, Si Cu, and Cr are selected to be about 0.5-l percent, 0.1-1.5 percent, 0.5-3.0 percent, 0.00-0.3 percent respectively of the total (Al+Mg+Si+ Cu+Cr) content. A typical alloy could contain 40 percent A1, 0.4 percent Mg, 0.25 percent Si, 0.! Cu, 0.1 percent Cr, balance Be impurities. This complex Be-Al-Mg-Si-Cu-Cr alloy can then be solution heat treated and precipitation hardened. It is about 20-150 percent stronger than its binary Be-Al counterpart, depending upon Be Content. The strength improvement persists to to at least 500 F.
EXAMPLE 15 The Be-Al-Mg-Si-Cu-Cr alloys described in Example 14 can be prepared in an alternate manner by first atomizing a melt of Be+ Al+ Si+ Cu+ Cr at approximately 2,500 F according to the teaching of U.S. Pat. No. 3,337,334. The Mg is then added as powder and the process carried out as in Example 14, with similar results.
EXAMPLE l6 Beryllium-aluminum base alloys containing -80 percent Al, 0.5-7.5 percent Zn, 0.2-4.0 percent Mg, 0.1-3.5 percent Cu, 0.0-0.3 percent Cr are prepared in a manner analogous to that described for a Be-Al-Mn alloy in Example 4. The amounts of Zn, Mg, Cu, and Cr are selected to be about 0.8-7.5 percent, 0.4-4.0 percent 0.2-3.5 percent 0.0-0.3 percent respectively of the total of (Al+Zn+Mg+Cu-+-Cr) contents. A typical alloy could contain 40 percent A1, 2.2 percent Zn, 1.0 percent Mg, 0.65 percent Cu, 0.1 percent Cr, balance Be impurities. The complex Be-Al-Zn-Mg-Cu-Cr alloy can be solution heat treated and precipitation hardened. Its strength is about 30-300 percent greater than that Be-Al the corresponding Be-Al binary alloy. This Be-Al-Zn-Mg-Cu alloy is also about 5-35 percent stronger than the Be-Al-Cu- Mg-Mn alloy described in Examples 12 and 13, but this advantage is lost at temperatures above about 200-250 F.
EXAMPLE 17 The Be-Al-Zn-Mg-Cu-Cr alloys described in Example 16, can be prepared in an alternate manner by first atomizing a melt of Be+Al+Cu+Cr at approximately 2,500 F according to the teaching of U. S. Pat. No. 3,337,334. The Mg and Zn are then added as powder(s) and the process carried out as in Example 16, with similar results.
EXAMPLE l8 Beryllium-aluminum base alloys containing 5-80 percent Al, 0.2-6.0 percent Cu, 0.2-5.0 percent Li, 0.1-1.0 percent Mn, 0.5-1.0 percent Cd are prepared in a manner analogous to that described for a Be-Al-Mn alloy in Example 4. The amounts of Cu, Li, Mn and Cd are selected to about 0.3-6 percent, 0.3-5 percent, 0.2-1.0 percent, and 0.1 to 1.0 percent respectively of the total of (Al+Cu+Li+Mn+Cd) contents. (It may be convenient to add Li as the hydride.) A typical alloy could contain 40 percent A], 1.8 percent Cu, 0.6 percent Li, 0.2 percent Mn, 0.1 percent Cd, balance Be impurities. This Be-Al-Cu-Li-Mn-Cd alloy can be solution heat treated and precipitation hardened. Its strength is about 30-300 percent greater than that of the corresponding Be-Al binary alloy. In addition the strength improvement persists to at least 600 F, and is superior to the Be-Al-Cu-Mg-Mn alloy described in Examples l2 and 13, throughout the temperature range of l00 to +600 F.
EXAMPLE 19 The Be-Al-Cu-Li-Cd-Mn alloy described in Example 18 can be prepared in an alternate manner by first atomizing a melt of Be+ Al+Cu+ Mn according to the teaching of U.S. Pat. No. 3,337,334. The Li and Cd are then added as powder(s) and the process carried out as in Example 13, with similar results. (It again may be convenient to add lithium as the hydride.)
EXAMPLE 20 -(See Example 4) is, in the solution-treated condition about 20 percent to 50 percent stronger than the binary Be-Al counter- P EXAMPLE 21 The alloy described in Example 20 can be prepared in an alternate manner by atomizing a melt of Be+Al+Ag at approximately 2,500 F according to the teachings of U. S. Pat. No. 3,337,334. Results are quite similar to those of Example 20.
If more complex alloys containing volatile addition elements (e.g., Zn or Mg) are desired, the Be+Al+Ag alloy powder may be mixed with volatile addition powders, and the mixture processed by the methods of Examples 4 and 5.
The preceding examples by no means exhaust the possibilities of this invention, but they do give an indication of what can be done and of the very significant results obtainable.
1. Beryllium-aluminum base alloys consisting essentially of 20 percent to percent by weight of beryllium with the remainder an aluminum-rich phase of aluminum containing in solid solution in the aluminum-rich phase at least one material from the group consisting of the elements magnesium, zinc, copper, lithium, silver, silicon and manganese, said material being present in an amount able to dissolve completely in said aluminum-rich phase alone at a temperature where the aluminum-rich phase is in a solid state, said alloy having a complexity index within the range of from approximately 1 to approximately 5, said beryllium being present as a network that is substantially continuous over large distances and said aluminum-rich phase being a single phase substantially continuous network separate from the beryllium network.
2. An alloy according to claim 1 having in addition in solid solution in the aluminum-rich phase at least one material from the group consisting of the elements titanium, chromium, bismuth, cadmium, lead, zirconium, iron, and nickel and in an amount able to dissolve completely in said aluminum-rich phase alone at a temperature where the aluminum-rich phase is in a solid state.
3. Beryllium-aluminum base alloys consisting essentially of at least 20 percent by weight beryllium with the remainder being primarily aluminum, which are characterized by a distinctive microstructural appearance in which the berylliumrich phase is present in the form of generally particulate, irregularly shaped substantially continuous networks which are interspersed by the aluminum-rich phase, which is a singlephase substantially continuous network, the mean free intercept distance, A, across the aluminum-rich phase being in a range of from approximately 1 to approximately 6 microns and in which the mean intercept diameter, D, across the beryllium-rich phase is in a range of from approximately 1 micron to approximately 16 microns, and at least one additional material chosen from the group consisting of the elements magnesium, zinc, copper, lithium, silver, silicon, and manganese, substantially all of said additional material being contained in solid solution in said aluminum-rich phase and being present in an amount no greater than what can be dissolved in said aluminum-rich phase at a temperature where the aluminum-rich phase is in a solid state.
4. An alloy according to claim 3 having in addition in solid solution in said aluminum-rich phase at least one material chosen from the group consisting of the elements titanium, chromium, bismuth, cadmium, lead, zirconium, iron, and nickel and being present in an amount no greater than what can be dissolved in said aluminum-rich phase at a temperature where the aluminum-rich phase is in a solid state.
5. An alloy composite consisting of beryllium-rich and aluminum-rich phases, each being a substantially continuous network, having a complexity index of approximately 1 to 5 and mean free path in the Be-rich or Al-rich phases of less than about 16 microns, in which the aluminum rich phase consists of a single-phase solid solution primarily of aluminum with at least one dissolved addition element selected from the group consisting of the elements in the following list, each said element being present in a concentration expressed as percent of the aluminum-rich constituent, no greater than Zinc up to 30 percent, Magnesium up to 17% percent, Copper up to 6 percent, Lithium up to 5% percent, Silver up to 10 percent, Silicon up to 2 percent, Manganese up to 1% percent,
the total beryllium concentration in the alloy composites being 20 to 90 percent of the whole.
6. An alloy composite according to claim 5 having in addition in solid solution in said aluminum-rich phase at least one element chosen from the group consisting of the elements in the following list in amounts as therein shown:
Titanium up to 1% percent, Chromium up to 1 percent, Bismuth up to 1 percent, Cadmium up to 1 percent, Lead up to percent, Zirconium up to a percent, lron up to 1/10 percent, Nickel up to 1/10 percent.
7. A ternary alloy composite consisting essentially of beryllium-rich and aluminum-rich phases, each being a single-phase substantially continuous network, and having a complexity index of approximately 1 to 5 and mean freepath in each of the Be-rich and Al-rich phases of less than about 16 microns, the aluminum-rich phase consisting essentially of an aluminum-magnesium solid solution with up to a maximum of about 17 r percent magnesium, the total beryllium concentration being to 95 percent of the whole.
8. A quaternary alloy composite consisting of 20 to 90 percent beryllium in a substantially continuous network and 10 to 80 percent of a single-phase aluminum-rich phase in a substantially continuous network containing in solid solution zinc and magnesium, the zinc content being up to percent of the aluminum-rich phase and the magnesium content being up to about 18 percent of the total aluminum-rich phase, the alloy composite having a complexity index of approximately 1 to 5.
9. A quaternary alloy composite according to claim 8 in which the aluminum-rich phase has been precipitation hardened.
10. A quaternary alloy composite consisting of 20 to 90 percent beryllium in a substantially continuous network and 10 to 80 percent of a single-phase aluminum-rich phase in a substantially continuous network containing in solid solution copper and magnesium, in which the copper content is up to 6 percent of the aluminum-rich phase and in which the magnesium content is up to 18 percent of the aluminum-rich phase, the alloy composite having a complexity index of approximately 1 to 5.
11. A quaternary alloy composite according to claim 10 in which the aluminum-rich phase has been precipitation hardened.
12. A complexed alloy composite consisting of 20 to 90 percent beryllium in a substantially continuous network, the remainder being an aluminum-rich phase in a single-phase continuous network containing in solid solution a plurality of addition elements selected from the group consisting of the elements in the following list at concentrations no greater than listed as a percent of total aluminum-rich phase, at least one of said addition elements being among the first seven on the list, said alloy composite having a complexity index of approximately 1 to 5:
Magnesium 17% Zinc 30 Copper 6 Lithium 5 A Silver 10 Silicon 2 Manganese 1% Titanium 1 9% Chromium 1 Bismuth 1 Cadmium 1 Lead 1% Zirconium b lron 1/1 0 Nickel l/ 10.
13. A method of making improved Be-Al base alloys which contain at least 20 percent by weight beryllium with the remainder being principally aluminum, comprising:
providing a Be-Al alloy powder containing 20 percent to percent beryllium, and having a Be-rich phase and an A]- rich phase, and characterized by a complexity index between about 1 and about 5, and
adding to the Be-Al alloy powder at least one addition material chosen from the group consisting of Mg, Zn, Cu, Li, Mn, Si, and Ag in an amount able to dissolve solely within the aluminum-rich phase, processing the resulting mixture to dissolve each such added addition material solely in the Al-rich phase forming a solid solution, said processing being done below the eutectic temperature of the combination of metals used,
whereby said complexity index remains substantially the same as before the addition material was added to the original Be-Al alloy.
14. The method of claim 13 wherein during said processing step, at least one material from the group consisting of Ti, Cr, Bi, Cd, Pb, Zr, Fe, and Ni is also added to the Al-rich phase under the same conditions forming a solid solution and in amount able to dissolve solely within the Al-rich phase at a temperature where the Al-rich phase is in a solid state.
15. The method of claim 13 wherein said processing step includes working the resultant blend at a temperature below the eutectic temperature of the alloy being worked to make the Al-rich phase more uniform and to improve the mechanical properties of the alloy.
16. The method of claim 13 wherein said processing step comprises solid state diffusion, including both strain-enhanced diffusion and surface difi'usion.
17. The method of making beryllium-aluminum alloys which contain at least 20 percent by weight beryllium with the remainder an aluminum-rich phase, being primarily aluminum but containing at least one other metal in solid solution, said alloys being characterized by their ability to undergo selective and substantial change in strength and ductility, said method comprising the steps of providing a starting material in the compositional range set forth above which is selected from the group consisting of fine aluminum coated beryllium powder and small particles of beryllium-aluminum alloy,
blending therewith without any melting at least one powdered addition product chosen from the group consisting of magnesium, zinc, copper, lithium, manganese, silicon, and silver, in an amount able to dissolve completely solely within said aluminum-rich phase at a temperature where the aluminum-rich phase is in a solid state,
compacting said particles under pressure, and
dissolving said addition product in solid solution in the aluminum-rich phase only of said beryllium-aluminum alloy, at a temperature below the eutectic temperatures of the blend and of the resulting alloy.
18. The method of claim 17 wherein said dissolving is accomplished by working the resultant blend at a temperature below the eutectic temperature of the alloy being worked, to enhance and hasten diffusion, to produce more uniformity in the aluminum-rich phase, and to improve the mechanical properties of the alloy.
19. The method of claim 17 wherein said dissolving step is done by solid state diffusion.
20. The method of claim 17 wherein in the blending step, at least one element from the group consisting of titanium, chromium, bismuth, cadmium, lead, zirconium, iron, and nickel is also blended in the same manner forming a solid solution and in amount able to dissolve completely solely within said aluminum-rich phase at a temperature where said aluminum-rich phase is in a solid state.
21. The method of claim 20 wherein said dissolving step comprises working the resultant blend at a temperature below the eutectic temperature of the alloy being worked to improve the uniformity and mechanical properties of the alloy.
22. The method of claim 20 wherein said dissolving step is done by solid state diffusion.
23. The method of making beryllium-aluminum base alloys of substantially improved strength and ductility and affording improved starting materials for more complex alloys, which contain at least 20 percent by weight beryllium with the remainder an aluminum-rich phase being primarily aluminum with at least one other component partitioned in solid solution to the aluminum-rich phase, which comprise the steps of melting a quantity of substantially pure beryllium, substantially pure aluminum and at least one material chosen from the group consisting of manganese, silicon, copper, silver, titanium, chromium, lead, zirconium, iron, and nickel, in the percentage range desired, said percentage range being no greater for each said material than the amount able to dissolve completely in said aluminum-rich phase alone, at a tempe ture where the aluminum-rich phase is in a solid state, atomizing said melt to form small particles of the desired alloy, rapidly quenching said particles, compacting said parti cles to form a substantially solid billet, outgassing said billet, heat treating said billet at an elevated temperature below the eutectic temperature of the alloy to form a solid solution of the said material in the aluminum-rich phase, and extruding the heated billet into the shape desired.
24. The method of claim 23 wherein these is added, after quenching and before compacting, material chosen from the group consisting of magnesium, zinc, lithium, bismuth, and cadmium, in an amount no greater than the amount able to dissolve completely within said aluminum-rich phase alone at a temperature where said aluminum-rich phase is in a solid state.
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|U.S. Classification||75/249, 420/542, 148/514, 420/401, 419/29, 148/415, 419/28, 148/419|
|International Classification||C22C21/00, C22C25/00, C22C1/00, C22C1/04|
|Cooperative Classification||C22C25/00, C22C1/0408, C22C1/0416|
|European Classification||C22C1/04B1, C22C25/00, C22C1/04B|