Abstract:
A 7XXX series aluminum alloy having reduced quench sensitivity suitable for use in aerospace structural components, such as integral wing spars, ribs, extrusions and forgings comprises, in weight %: 6 to 10 Zn, 1.3 to 1.9 Mg, 1.4 to 2.2 Cu, wherein Mg≦Cu+0.3, one or more of 0 to 0.4 Zr, up to 0.4 Sc, up to 0.2 Hf, up to 0.4 Cr, up to 1.0 Mn and the balance Al plus incidental additions including Si, Fe, Ti and the like plus impurities. By controlling the Mg content to 1.3 to 1.7 wt. %, limiting Mg≦Cu+0.3 and 6.5≦Zn≦8.5, the alloy provides significantly improved combined strength and fracture toughness in heavy gauges. For example, in a six-inch thick plate there is provided a combination of about 75 ksi quarter-plane tensile yield strength (L) with a fracture toughness (L−T) of about 33 ksi{square root}in which progresses by artificial aging/tempering to a combined strength and fracture toughness of about 67 ksi tensile yield strength (L) and a fracture toughness (L−T) of about 40 ksi{square root}in. The alloy product possesses equally attractive combinations of strength and fracture toughness when intentionally quenched slowly following solution heat treatment so as to lessen dimensional distortion, particularly in shapes of varying cross section.

Description:
BACKGROUND OF THE INVENTION  
         [0001]    1. Field of the Invention  
           [0002]    The present invention relates generally to aluminum alloys and, more particularly, to 7XXX series aluminum alloys having superior strength-toughness combinations suitable for thick gauge (e.g., 2-10 inches) structural parts in aerospace applications. Parts made from the alloy of the present invention find specific utility as structural components such as integral spar members and the like which are integrally machined from thick sections, including rolled plate, used in the construction of structural members for high capacity aircraft. The alloy may also be formed by other known hot forming techniques such as extrusion and forging. The forging process is particularly suitable for manufacturing high strength aircraft components, such as, for example, main landing gear beams.  
           [0003]    2. Background of the Invention  
           [0004]    As the size of new aircraft, or modification of current models, gets larger to accommodate heavier payload and/or longer flight range for improved performance and economy, the demand on weight savings of structural components, such as fuselage, wing, spar, etc., continues to increase. The aircraft industry is meeting this demand through specification of higher strength requirements of metal parts to enable reduced section thickness as a weight savings expedient. However, during the late 1960&#39;s, those skilled in the field of fracture mechanics discovered that, in addition to the strength, the durability and damage tolerance of the material are also critical to the fail-safe aircraft structural design. Such consideration of multiple material attributes for aircraft applications eventually led to today&#39;s damage tolerance design, which combines principles of fail-safe design with periodic inspection techniques.  
           [0005]    Traditionally, an aircraft wing structure comprises a wing box designated generally by reference numeral  2  in FIG. 1. The wing box  2  is attached to upper and lower wing skins  4  and  6 , respectively, connected by vertical structural members called spars  12  and  20 . During flight, the upper wing structure of a commercial aircraft is compressively loaded, calling for high strength with acceptable fracture toughness as desired material attributes. The upper wing is typically built with a 7XXX series aluminum alloy such as 7150 (U.S. Reissue Pat. No. 34,008) or 7055 (U.S. Pat. No. 5,221,377). The lower wing structure of a commercial aircraft is under tension during flight and, therefore, requires high damage tolerance. Although it is desirable to design the lower wing with a high strength alloy to maximize weight efficiency, the damage tolerance characteristics of the high strength 7XXX series alloys fall short of the need. Therefore, most commercial aircraft manufacturers specify a damage-tolerant 2XXX series aluminum alloy such as 2024 or 2324 for lower wing applications (U.S. Pat. No. 4,294,625), both of which are significantly lower in strength than the 7XXX series upper wing alloys. The alloy members and temper designations used herein are in accordance with the well-known aluminum alloy product standards of the Aluminum Association.  
           [0006]    The upper and lower wing skins  4  and  6 , respectively, are typically stiffened by longitudinally extending stringer members  8 ,  10  in J, I, T, Z or other shapes which are fastened to the inside surfaces of the wing skins, as shown in FIG. 1. The upper wing stringers  8  and upper spar caps  14 ,  22  are usually manufactured from a 7XXX series aluminum alloy, and the lower wing stringers  10  and spar caps  16 ,  24  with a 2XXX series aluminum alloy for the same structural reasons discussed above regarding their respective strength and damage-tolerance properties. Vertical web members  18 ,  26  fastened to both the upper and the lower spar caps run along the longitudinal direction of the wing, constituting the structural member spars  12  and  20 . This traditional spar design is also called a “built-up” spar, comprising the upper spar cap  14  or  22 , the web  18  or  20  and the lower spar cap  16  or  24  with fasteners (not shown). Obviously, the fasteners and fastener holes at the joints are structural weak links. In order to ensure the structural integrity of the built-up spars  18  and  20 , many parts of the web and spar caps have to be thickened, thereby adding weight to the structure.  
           [0007]    A desirable design approach to overcome the aforementioned weight penalty imposed by the built-up spar is to make the upper spar, the web and the lower spar out of a single piece of aluminum alloy product. This design eliminates the web-to-upper-spar and web-to-lower-spar joints, and can provide significant weight savings with respect to the built-up spar. The one-piece spar is called an “integral spar,” and can be machined to final shape from a thick plate, a thick extrusion or a thick forging. The integral spar not only weighs less, but also involves less cost to manufacture and assemble due to the elimination of fasteners. The ideal alloy for the integral spar is one which has the same strength as the upper wing skin with sufficient fracture toughness to meet the damage tolerant requirements of the lower skin. However, current alloys do not meet the desired property requirements. For example, the lower strength of alloy 2024-T351, commonly used in lower wing applications, will not be able to safely carry the load transmitted from the highly loaded upper wing unless its section thickness is significantly increased, which adds undesirable weight to the aircraft. Alternatively, to design the upper wing according to the 2XXX strength capability would result in overall weight penalty. The higher strength 7XXX series-T7X alloys in thick gauges, such as 7050-T74 (U.S. Pat. No. 3,881,966) still fall short of the need in strength in the required thickness, as elaborated below.  
           [0008]    Emerging large jet aircrafts require very large wings which, in turn, require integral spar alloy products as thick as 6 to 8 inches or more for weight efficiency. Today&#39;s industry standards for thick plate, for example, 6 inch thick 7050-T7451, are given in the Aerospace Materials Specifications AMS 4050F, setting forth a minimum yield strength of 60 ksi in the longitudinal (L) direction and plane strain fracture toughness, K Ic  (L−T) of 24 ksi{square root}in, and those for the transverse direction (LT and L-T) are, respectively, 60 ksi and 22 ksi{square root}in. However, the most recently developed upper wing skin specified on the Boeing 777, e.g., aluminum alloy 7055-T7751, is capable of meeting a yield strength minimum of 86 ksi according to the MIL-HDBK-5H. Therefore, if an integral spar with 60 ksi yield strength is used, the strength capability of the upper wing skin will not be able to be taken full advantage of for maximum weight efficiency. Hence, a higher strength, thick spar alloy with sufficient fracture toughness is highly desirable for the construction of an integral spar. This is but one specific example of the benefits of an aluminum material with high strength and toughness in thick sections, but many others exist in modern aircraft.  
           [0009]    It is known that the various tempers resulting from different artificial aging treatments will impart different levels of strength and other performance characteristics such as corrosion resistance and fracture toughness. The 7XXX series aluminum alloys are nominally in artificially aged tempers such as peak strength T6-type or over-aged T7-type tempers. U.S. Pat. Nos. 4,863,528; 4,832,758; and 4,477,292, along with U.S. Pat. No. 5,108,520 all describe tempers for 7XXX series alloys to provide a range of combinations of strength and performance. All of the aforesaid patents are fully incorporated herein by reference. It is well-known to those skilled in the art that for a given 7XXX series alloy, the peak strength T6-type temper provides the highest strength combined with the lowest fracture toughness and corrosion resistance. It is also known that the most over-aged temper such as T73-type temper for the same alloy provides the highest fracture toughness and corrosion resistance combined with the lowest strength. In practice, an appropriate temper is chosen somewhere between these two extremes to suit a particular application. A more complete description of the temper description (the “T-XX” suffix) is given in the Aluminum Standards and Data 2000, published by The Aluminum Association, Inc.  
           [0010]    The processing of most aerospace alloys generally requires a solution heat treatment (SHT) followed by quenching and subsequent artificial aging. In the quest for a stronger material, alloy designers are faced with two natural phenomena. As the product shape gets thicker, the quench rate experienced at the interior of the product cross section naturally decreases, which results in a loss of strength and fracture toughness in the product. This important phenomenon is known as “quench sensitivity”. Also, there is an inverse relationship between strength and fracture toughness; that is, as the alloy gets stronger, its toughness tends to decrease. Up until now, no one has been able to solve the problem concerning how to avoid these conflicting phenomena and provide a stronger, thick alloy product combined with appropriate high fracture toughness.  
           [0011]    To better understand the present invention, it is helpful to note certain demonstrated trends in the art concerning 7XXX series aluminum alloy development efforts to meet demands for improved properties.  
           [0012]    Type 7050 aluminum alloy employs Zr in place of Cr as a dispersoid agent and realizes a significant improvement in quench sensitivity over the prior 7075 alloy and, thus, has been the mainstay for thick section application as plate, extrusion or forging in aerospace service. To meet higher strength requirements, the alloy contents in 7050 alloy are higher than in 7075. For example, the lower limit of Cu (composition as registered with the Aluminum Association) increased from 1.2 weight % in 7075 to 2.0 weight % in 7050 alloy and, likewise, Zn from 5.1 to 5.7 weight %. In 7150 alloy for upper wing application with still higher strength-toughness requirements, the Mg and Zn lower limits also slightly increased compared to those in 7050, namely, Zn increased from 5.7 to 5.9 weight %, and Mg from 1.9 to 2.0 weight %, respectively, while simultaneously the metal purity was also improved by lowering the Fe and Si levels. A newer generation upper wing skin alloy, represented by 7055 alloy, provides a 10 percent improvement in compression yield strength properties and employs compositions with a much higher Zn range of 7.6 to 8.4 weight % while maintaining similar Cu levels and a slightly lower Mg range at 1.8 to 2.3 weight % compared to either 7050 or 7150 alloys.  
           [0013]    Thus, the past trend against any drive for higher strength had been to increase alloying additions and composition optimization as dictated by phase equilibrium relations, while metal purity increase and microstructure control through thermal-mechanical processing (TMP) were also sought in order to obtain simultaneous improvements in toughness and fatigue life, amongst other properties.  
           [0014]    U.S. Pat. No. 5,865,911 reports a significant improvement in toughness at equivalent strengths in plates of a 7XXX series alloy less than 2.5 inches thick. The improvement was obtained through optimization of alloy compositions by careful use of phase equilibrium relations, coupled with other innovations. However, the purported superior strength-toughness advantage was significantly reduced in this alloy when a quench rate simulating the mid-plane of a 6-inch product is used after SHT, showing only limited improvement in properties over that of 7050 alloy. Thus, quench sensitivity in this alloy has been responsible for the loss in property advantages in thicker gauges.  
           [0015]    In another example, 7040 alloy registered with the Aluminum Association reports the following composition, in weight percent, for the major alloying elements: 5.7 to 6.7 Zn, 1.7 to 2.4 Mg and 1.5 to 2.3 Cu. Published literature (Shahani et al., “High Strength 7XXX Alloys For Ultra-Thick Aerospace Plate: Optimization of Alloy Composition,” Proc. ICAA-6, Vol. 2, 1998; see, also, U.S. Pat. No. 6,027,582 to Shahani et al.) teaches that an optimization balance was pursued between alloying additions in an effort to improve strength and other properties while avoiding excess additions to minimize quench sensitivity. While some property improvements over 7050 alloy were claimed in thicker gauges, the improvements of 7040 still fell short of those desired for newer commercial aircraft designs.  
           [0016]    The instant invention solves the problems encountered in the prior art by providing a 7XXX series aluminum alloy which exhibits significantly reduced quench sensitivity so as to provide significantly higher strength and fracture toughness levels than heretofore possible in thick gauge aerospace structural members.  
         SUMMARY OF THE INVENTION  
         [0017]    Briefly stated, the present invention is directed to a 7XXX series aluminum alloy having significantly reduced quench sensitivity in thick gauges, i.e., greater than about 2 inches and, more preferably, in the thickness range of about 4 to 8 inches or greater. A presently preferred broad composition of the alloy of the present invention consists essentially of, in weight %: about 6 to 10 zinc (Zn), about 1.3 to 1.9 magnesium (Mg), about 1.4 to 2.2 copper (Cu), wherein Mg≦Cu+0.3, about 0 to 0.4 zirconium (Zr), about 0 to 0.4 scandium (Sc), about 0 to 0.2 hafnium (Hf), about 0 to 0.4 chromium (Cr), about 0 to 1.0 manganese (Mn), the balance being aluminum (Al) and other incidental elements.  
           [0018]    A more narrow, presently preferred alloy composition according to the present invention consists essentially of, in weight %: about 6 to 8.5 Zn, about 1.3 to 1.8 Mg, about 1.4 to 2.0 Cu, wherein, Mg≦Cu+0.3, one or more elements selected from the group consisting of up to about 0.4 Zr, up to about 0.4 Sc, up to about 0.2 Hf. up to about 0.4 Cr, up to about 1.0 Mn, the balance being Al , incidental additions and impurities.  
           [0019]    A still more narrowly defined, presently preferred alloy composition according to the present invention consists essentially of, in weight %: about 6.5 to 8.5 Zn, about 1.3 to 1.7 Mg, about 1.4 to 2.0 Cu, wherein Mg≦1.7 and Mg≦Cu+0.3, one or more elements selected from the group consisting of up to about 0.4 Zr, up to about 0.38 Sc, up to about 0.20 Hf. up to about 0.37 Cr, up to about 1.0 Mn, the balance being Al and other incidental additions and impurities. The above defined alloys may contain impurities and other incidental/intentionally made additions common and well-known to the 7XXX series family of aluminum alloys, such as on the order of, for example, in weight %: nominally about 0.03 or up to about 0.12 maximum silicon (Si), nominally about 0.05 or up to about 0.15 max. iron (Fe), nominally about 0.025 or up to about 0.15 max. titanium (Ti) and the like. The “other” additions are generally governed by the 0.05-0.15 ranges as defined in the alloy designations by the Aluminum Association.  
           [0020]    The elements Zr, Sc, Hf. Cr and Mn are introduced as dispersoid forming elements aimed at providing an unrecrystallized or partially recrystallized grain structure in the invention wrought products. Such grain structures are required to achieve the highest combination of strength, fracture toughness and stress corrosion resistance. The above described dispersoid forming elements are substantially in supersaturation after casting, and form fine dispersoid particles by solid state reactions during thermal-mechanical processing. Zr forms Al 3 Zr, Sc forms Al 3 Sc, Hf forms Al 3 Hf, Cr forms either Al 12 Mg 32 Cr or Al 18 Mg 3 Cr 2  and Mn forms Al 20 Cu 2 Mn 3  dispersoid particles. The dispersoid particles retard or stop recrystallization by exerting a drag force on the recrystallization nuclei. The homogenization process for the invention alloys, in fact, is aimed not only at evenly redistributing and dissolving the cored micro-segregation of the major alloying elements but also at causing an optimum, copious distribution of dispersoids for controlling the grain structure during subsequent thermal-mechanical processing. When multiple dispersoid elements are present, there may be synergistic effects for grain structure control. The optimum dispersoid content depends, in part, on the solidification process and, in part, on the thermal-mechanical process.  
           [0021]    The alloys of the present invention are conventionally prepared by melting and may be direct chill (D.C.) cast into ingot form. Conventional grain refiners such as titanium boride may also be used as well-known in the art. After conventional scalping and homogenization (if needed), the ingots are further processed by, for example, hot rolling into plate or extrusion or forging into special shaped sections. Generally, the heavy sections are on the order of greater than 2 inches and, more typically, on the order of 4, 6, 8 or up to 12 inches or more in thickness. In the case of heavy plate of about 4 to 8 inches in thickness, the plate is solution heat treated and quenched and mechanically stress relieved such as by stretching, for example, up to 8%, or compression. A desired structural shape is then machined from these heat treated heavy plate sections to form the final part, such as, for example, an integral wing spar. Similar SHT, quench and often stress relief operations are also followed in the manufacture of thick sections made from extrusions and forgings.  
           [0022]    Good combinations of properties are desired in all thicknesses, but they are particularly useful in thickness ranges where, conventionally, as the thickness increases, quench sensitivity of the product also increases. Hence, the alloy of the present invention finds particular utility in heavy gauges of, for example, greater than 4 inches in thickness up to 12 inches or more. 
       
    
    
     BRIEF DESCRIPTION OF THE DRAWINGS  
       [0023]    [0023]FIG. 1 is a transverse cross-sectional view of a typical wing box construction of an aircraft including front and rear spars of conventional three-piece built-up design;  
         [0024]    [0024]FIG. 2 is a graph showing two calculated cooling curves to approximate the mid-plane cooling rates for 6- and 8-inch thick plates under spray quenching, and superimposed on top are two experimental cooling curves simulating the cooling rates of a 6-inch thick and an 8-inch thick plate;  
         [0025]    [0025]FIG. 3 is a graph showing tensile yield strength TYS (L) versus fracture toughness K q  (L-T) relations for selected alloys of the present invention and other alloys including type 7150 and 7055 aluminum alloys as controls, all based on simulation of quarter-plane quench rates of 6-inch thick plate, extrusion or forging;  
         [0026]    [0026]FIG. 4 is a graph similar to FIG. 3 showing tensile yield strength TYS (L) versus fracture toughness K q  (L-T) relations for selected alloys of the present invention and other alloys including 7150 and 7055 controls, all based on simulation of quarter-plane quench rates of 8-inch thick plate, extrusion or forging;  
         [0027]    [0027]FIG. 5 is a graph showing the influence of Zn content on quench sensitivity as demonstrated by directional arrows for the tensile yield strength changes in a 6-inch thick plate quench simulation;  
         [0028]    [0028]FIG. 6 is a graph showing the influence of Zn content on quench sensitivity as demonstrated by directional arrows for the tensile yield strength changes in an 8-inch thick plate quench simulation;  
         [0029]    [0029]FIG. 7 is a graph showing cross plots of tensile yield strength TYS (L) versus plane-strain fracture toughness K Ic ,(L−T) values at quarter plane of a full-scale production 6-inch thick plate of the invention alloy and of 7050 and 7040 alloys from literature; and  
         [0030]    [0030]FIG. 8 is a graph showing the influence of section thickness on tensile yield strength, as an index of quench sensitivity property, from a full-scale production die-forging study comparing alloys of the invention and 7050 alloys. 
     
    
     DETAILED DESCRIPTION OF THE INVENTION  
       [0031]    Mechanical properties of importance for the thick plate, extrusion or forging for aircraft structural products, as well as other non-aircraft structural applications, include strength, both in compression as for the upper wing skin and in tension for the lower wing skin. Also of importance are the fracture toughness, both plane-strain and plane-stress, and corrosion resistance such as, for example, exfoliation and stress corrosion cracking resistance.  
         [0032]    Integral wing spars and wing skin panels with integral stringers must be machined from relatively thick plates or other structural shapes which are extruded or forged from thick ingots or billets and which have been then solution heat treated, quenched and artificially aged. It is not always feasible to solution heat treat and quench the finished structural products because the rapid cooling of the quenching step would induce residual stress and cause dimensional distortion. The quench-induced residual stress could also possibly cause stress corrosion cracking, and re-work to straighten parts associated with dimensional distortion could render assembly impracticably difficult. While it is much easier to obtain better mechanical properties in thinner cross sections because of the faster cooling which prevents unwanted precipitation of alloying elements, this cannot be done when quench distortion is present. Hence, it is necessary to solution heat treat and quench the thick plates, extrusions or forgings, mechanically straighten and flatten the wrought product while simultaneously relieving the residual stress, followed by artificial aging to produce the desired final temper. The wrought product is then machined to achieve the desired shape of the finished structural component.  
         [0033]    As alluded to above, in solution heat treating and quenching thick sections, the quench sensitivity of the aluminum alloy is of great concern. After solution heat treating, it is desirable to quickly cool the material in order to retain the various alloying elements in solid solution rather than to allow them to precipitate out of solution in coarse form as occurs in slow cooling. The latter occurrence producing the coarse precipitates results in a decline in mechanical properties. In thick product cross sections, i.e., over 3 inches, and more particularly in heavier sections of 4 to 8 inches or more, the quenching medium acting on the exterior surfaces of the workpiece (such as a plate, forging or extrusion, for example) cannot efficiently extract heat from the center or mid-plane region of the material. This is due to the physical distance to the surface and the fact that heat is extracted through the metal by conduction which is distance dependent. In thinner product cross sections, quench rate at the mid-lane is naturally higher than that in heavier sections. Hence, the quench sensitivity property of an alloy is not as important in thinner gauge shapes as it is in heavier gauge workpieces.  
         [0034]    The present invention is, of course, directed at increasing the strength-toughness properties in a 7XXX series aluminum alloy in thicker gauges, i.e., greater than about 1.5 inches, and, thus, the quench sensitivity of the alloy is of extreme importance. In thicker gauges, the less quench sensitivity the better, with respect to a material&#39;s ability to retain alloying elements in solid solution (to avoid the formation of coarse precipitates upon slow cooling) particularly in the slowest-cooled mid-plane region. The present invention achieves the goal of lower quench sensitivity by providing a carefully controlled alloy composition which permits the manufacture of heavier gauges while achieving superior strength-toughness properties.  
         [0035]    The present invention resulted from an exploration for possible property advantages in 7XXX series alloys based upon certain phase equilibrium features that had been identified by the present inventors in the Al—Zn—Cu—Mg alloy system. The intent was to explore if increased additions of Zn could expand the matrix phase field (based on aluminum with a face centered cubic crystal structure) to allow for increased solubility of other alloying additions such as Cu and Mg for increased strength.  
         [0036]    The experimental work consisted of the following: Twenty-eight 11-inch diameter ingots were direct chill (DC) cast, homogenized and extruded into 1.25 inch thick by 4-inch wide rectangular bars. The bars were next solution heat treated and were quenched at different rates to mimic cooling conditions for thin section as well as those approximating 6-inch and 8-inch thick sections. The bars were then cold stretched by 1.5 percent for residual stress relief. The compositions of alloys studied are set forth in Table 1 below, in which the Zn content ranged from 6.0 weight % to slightly in excess of 11.0 weight %, while Cu and Mg each were varied between 1.5 and 2.3 weight %, respectively.  
                                                                                                                       TABLE 1                               Invention   Composition       Invention   Composition       Specimen   Alloy   (wt. %)   Specimen   Alloy   (wt. %)            No.   Y/N   Cu   Mg   Zn   No.   Y/N   Cu   Mg   Zn                    57   Y   1.57   1.55   6.01   71   N   1.86   1.93   10.93       58   N   1.64   2.29   5.99   72   N   1.98   2.09   11.28       59   N   2.45   1.53   5.86   73   N   1.97   1.86   9.04       60   N   2.43   2.26   6.04   74   Y   1.48   1.50   9.42       61   N   1.95   1.94   6.79   75   N   1.75   2.29   9.89       62   Y   1.57   1.51   7.56   76   N   2.48   1.52   9.60       63   N   1.59   2.30   7.70   77   N   2.19   2.19   9.74       64   N   2.45   1.54   7.71   78   N   1.68   1.55   11.38       65   N   2.46   2.31   7.70   79   N   1.65   2.28   11.04       66   N   2.05   1.92   8.17   80   N   2.38   1.53   11.08       67   Y   1.53   1.52   8.65   81   N   2.22   1.97   9.04       68   N   1.57   2.35   8.62   82   N   1.79   2.00   10.17       69   N   2.32   1.45   8.25   83   N   2.23   2.28   6.62       70   N   2.04   2.19   8.33   84   N   2.48   1.98   8.31                  
 
         [0037]    For all alloys other than the controls: Target Si=0.03, Fe=0.05, Zr=0.12, Ti=0.025  
         [0038]    For 7150 Control (S# 83): Target Si=0.05, Fe=0.10, Zr=0.12, Ti=0.025  
         [0039]    For 7055 Control (S# 84): Target Si=0.07, Fe=0.11, Zr=0.12, Ti=0.025  
         [0040]    Different quenching approaches were explored in order to obtain at the mid-plane of the 1.25 inch thick extruded bars a cooling rate mimicking that at the mid-plane of a 6-inch thick plate subjected to spray quenching in 75° F water in a full-scale production plant. A second simulation involved mimicking, under identical circumstances, a bar cooling rate corresponding to that of an 8-inch thick plate.  
         [0041]    The quenching simulation involved modification of the heat transfer characteristics of the quenching medium, as well as the surface of the part by immersion quenching of the extruded bars with simultaneous incorporation of three practices: a defined warm water temperature, saturation of the water with CO 2  gas, d a chemical treatment of the bars to render a bright etch surface finish to it.  
         [0042]    For mimicking the 6-inch thick plate cooling condition, the water temperature for immersion quenching was held at 180° F., the solubility level of CO 2  in the water was about 0.20 lan (a measure of dissolved CO 2  concentration, lan=Standard volume of CO 2 /Volume of water), and the surface had a standard bright etch finish.  
         [0043]    For the 8-inch thick plate cooling simulation, the water temperature utilized was 190° F. with a CO 2  solubility reading varying between 0.17 and 0.20 Ian together with a standard bright etch surface finish.  
         [0044]    The cooling rates were measured by thermocouples inserted into the mid-plane of the bar samples. For benchmark reference, the two calculated cooling curves to approximate the mid-plane cooling rates under spray quenching at the plants for the 6-and 8-inch thick plates were plotted, FIG. 2. Superimposed on them were displayed two groups of plots, the lower group (in the temperature scale) representing cooling curves simulating mid-plane of the 6-inch and the upper group simulating mid-plane of the 8-inch thick plate cooling rates. The cooling rates of the simulation were very similar to those of the plant production plates in the important temperature range above 500° F., although the simulated cooling curves for the experimental material differed from those of the plant plates below 500° F., which was not considered critical. Hence, the simulation was considered successful.  
         [0045]    After solution heat treating and quenching, aging behaviors were studied using multiple aging times and aided by EXCO (exfoliation corrosion resistance), EC (electrical conductivity) and SCC (stress corrosion cracking resistance) measurements. A two-step aging practice was used that consisted of a slow heat-up in 6 hours to 250° F., a 4-hour soak at 250° F. followed by second step aging soak at 320° F. for variable times ranging from 4 to 36 hours depending on the alloy composition in order to obtain acceptable EC and EXCO readings. The different aging steps were optimized not only to obtain improved strength-toughness properties, but also to obtain good corrosion resistance.  
         [0046]    Tensile and compact tension plane-strain fracture toughness test data were collected on samples that had been given the different minimum aging times required in order to obtain a visual EXCO rating of EB or better (EA) for an acceptable exfoliation corrosion resistance property, and an electrical conductivity (EC) reading minimum of 36% IACS (International Annealed Copper Standard), also used as an indicator for degree of over-aging and corrosion resistance. All tensile tests were performed according to the ASTM specification E8, and all plane-strain fracture toughness according to the ASTM specification E399.  
         [0047]    The original intent in the investigation, as stated above, was to use increased solutes in an expanded alloy phase field, made possible through high Zn additions, in an effort to obtain improved strength-toughness properties. However, it was surprising to find that, contrary to our expectations, the best combinations of strength-toughness properties were obtained at much lower solute compositions of Cu and Mg than anticipated or heretofore tried in the industry.  
         [0048]    We have found that it is desirable to use Zn at higher levels (above 6.7 wt. %, preferably around 7.5 wt. %), in contrast with the recent 7040 alloy with Zn specified between 5.7 and 6.7 wt. %, as well as the earlier 7050 and 7010 alloys.  
         [0049]    We have also found that for the best strength-toughness properties, a concomitant decrease in the Cu and Mg levels is required. Thus, the Zn levels in the alloy of our invention are higher, and the Cu and Mg levels are lower than those specified and practiced in the prior art. The Mg range for 7040 alloy is 1.7 to 2.4 wt. %. The Mg range for 7050 alloy is 1.9 to 2.6 wt. % and for 7010 alloy 2.1 to 2.6 wt. %. The desirable Mg range in the present invention is 1.3 to 1.9 wt. %, with a more preferred range being 1.3 to 1.7 wt. % for best properties.  
         [0050]    The Cu range registered by the Aluminum Association for 7040 is 1.5 to 2.3 wt. %, while the Cu range for the present invention is 1.4 to 2.2 wt. %, with a more preferred content at no more than 2.0 wt. % and, still more preferably, no more than 1.8 wt. % Cu for best properties. We have also discovered that the performance is enhanced when Cu is greater than Mg, or when Cu+0.3 wt. % is greater than or equal to Mg, contrary to the disclosure of 7040 (U.S. Pat. No. 6,027,582), where Mg content is required to be higher than Cu.  
         [0051]    [0051]FIG. 3 shows the strength-toughness plot of results from slow quenching of alloy samples of Table 1 from the SHT temperature to simulate a 6-inch thick product. As is readily noticed in these plots, a family of alloy compositions stood out, namely, alloy numbers 57, 62, 67 and 74 (upper part of the diagram) from the rest which displayed very high fracture toughness combined with high strength properties. Surprisingly, all these alloy compositions belonged to the low-Cu and low-Mg composition ends of our choice, namely, at around 1.5 wt. % Mg together with 1.5 wt. % Cu, while Zn levels varied from 6.0 to 9.5 wt. %.  
         [0052]    The Zn levels for the improved alloys illustrated were about: 6.0 wt. % for alloy #57, 7.5 wt. % for alloy #62, 8.6 wt. % for alloy #67 and 9.5 wt. % for alloy #74.  
         [0053]    The substantial improvement in the strength-toughness properties can be seen by comparison against the two control alloys 7150 (alloy #83) and 7055 (alloy #84) which had been processed in identical manner. In FIG. 3, a dotted line connecting the two data points has been drawn to show their strength-toughness property trend which positions the control alloys considerably below the low solute data points discussed.  
         [0054]    Also included in the plots are results for alloys with about 1.9 wt. % Mg and 2.0 wt. % Cu at varying Zn levels of 6.8 wt. % (alloy #61), 8.2 wt. % (alloy #66), 9.0 wt. % (alloy #73) and 10.2 wt. % (alloy #82). The results show a dramatic drop in toughness in these alloys compared to the 1.5 wt. % Mg and 1.5 wt. % Cu containing alloys at each of the corresponding Zn levels. The strength-toughness properties in thick gauge product in these higher Mg and Cu alloys are similar to or marginally better than the 7150 and 7055 control alloy trend line. The results clearly demonstrate the strong degradation in properties that occurs with a moderate increase in Cu and Mg above the low Cu and Mg levels in the alloys of the present invention, and approaching the levels of many of the current commercial alloys.  
         [0055]    A similar set of results with similar conclusions as above is depicted in FIG. 4 for a quench condition which was even slower than in FIG. 3. The condition of FIG. 4 roughly approximates an 8-inch thick plate mid-plane cooling condition.  
         [0056]    Thus, unlike the past teachings, some of the highest strength-toughness properties were obtained at some of the leanest Cu and Mg levels not considered thus far in the current commercial alloys. Concomitantly, the Zn levels at which the properties were most optimized corresponded to levels much higher than those cited in type 7040, 7050 or 7010 aluminum thick product alloys.  
         [0057]    The substantial improvement in strength-toughness properties in quench sensitive thick sections in the alloys of the invention conceivably resulted from a significant reduction in the quench sensitivity of the alloys owing to the lean Cu and Mg compositions, while higher Zn additions facilitated this through expanding the stability of the matrix phase field. This conclusion is supported in FIG. 5 where the strength (TYS) increased gradually with increasing Zn content for alloys #57 to #62 to #67. Thus, unlike the past teachings in the literature, including that for 7040 alloy, higher solute in the case of Zn did not hurt quench sensitivity and was, in fact, proven to be beneficial against slow quench conditions, i.e., thick sections. However, at a still higher Zn level of 9.4 wt. % the resultant enlarged phase field could not further compensate for the increased solute level, and the strength (TYS) of alloy #74 dropped below the other lower Zn alloys above, see FIG. 5. Also, with further, still slower quench conditions in simulated 8-inch thicknesses as depicted in FIG. 6, the quench sensitivity increased even at 8.7 wt. % Zn as depicted by the strength (TYS) of alloy #67 displaced below that of alloy #62 with 7.6 wt. % Zn. The high solute effect on quench sensitivity is also depicted in the relative positioning on the strength axis of the two control alloys, 7150 (alloy #83) and 7055 (alloy #84). 7055 was stronger than 7150 under slow quench (FIG. 5), but the relative scale was reversed under still slower quench conditions as in FIG. 6.  
         [0058]    It is also of interest to note alloy #63, which, according to Table 1, contained 1.59 wt. % Cu, 2.30 wt. % Mg and 7.70 wt. % Zn, (i.e., the Mg&gt;Cu). Alloy #63 exhibited in FIG. 3 a high strength (TYS) of about 73 ksi but a relatively low fracture toughness K q  (L−T) of about 23 ksi{square root}in. By way of comparison, alloy #62, which contained 1.57 wt. % Cu, 1.51 wt. % Mg and 7.56 wt. % Zn (i.e., Mg&lt;Cu), exhibited in FIG. 3 a strength (TYS) of greater than 75 ksi and a high fracture toughness of about 34 ksi{square root}in (a 48% increase in toughness). The data show the criticality of maintaining the Mg content below 1.7 wt. % as well as the importance of keeping the Mg content generally and preferably no greater than the Cu content or no grater than Cu+0.3 wt. %.  
         [0059]    It is desirable to achieve optimum and/or balanced fracture toughness (K q ) and strength (TYS) properties in the alloys of the present invention. As can be seen and appreciated by comparing the compositions set forth in Table 1 with the corresponding fracture toughness and strength values plotted in FIG. 3, those sample alloys falling within the compositions of the present invention achieve this balance of properties. It will be seen that the alloy samples having compositions according to the present invention, namely, alloys #57, #62, #67, and #74, either possess a fracture toughness value (K q ) in excess of about 34 ksi{square root}in with a tensile yield strength (TYS) of greater than 69 ksi or they possess a fracture toughness value of greater than about 29 ksi{square root}in combined with a higher tensile yield strength of greater than about 75 ksi.  
         [0060]    It will also be seen that the upper limit of the Zn content is important in achieving balanced toughness and strength properties. Those alloy samples which exceeded the upper limit on Zn of 11.0 wt. %, such as alloy #80 (11.08 wt. % Zn), or alloy #78 (11.38 wt. % Zn), failed to possess the minimum combined strength and fracture toughness levels set forth above for alloys of the invention.  
         [0061]    The alloy of the present invention, thus, provides high damage tolerance in thick aerospace structures resulting from its enhanced, combined fracture toughness and yield strength properties.  
         [0062]    K q  values are results of plane strain fracture toughness tests that do not conform to validity criteria of ASTM E399. In the current tests yielding K q  values, the validity criteria that were violated were: (1) Pmax/Pq≦1.1 primarily, and (2) B (thickness)≧2.5 (K q /σ ys ) 2  occasionally, where K q , σ ys , P max , and P q  are as defined in ASTM E399-90. These invalidities are a consequence of the high fracture toughness of the invention alloy. A thicker and wider specimen than that permitted by the extruded bar (1.25 inch thick×4 inch wide) would have been required to obtain valid plane-strain K Ic  results. A valid K Ic  is generally considered a material property relatively independent of specimen size and geometry. K q  on the other hand is not a material property in the strict sense because it can vary with specimen size and geometry. However, typically, K q  values from specimens smaller than that needed to obtain a valid result are conservative with respect to K Ic , in other words, the reported fracture toughness (K q ) values are, in fact, generally lower than the standard K Ic  values obtained when sample size related validity criteria of ASTM E399-90 are satisfied. The values of K q  were obtained using compact tension test specimen per ASTM E399 with a thickness, B, of 1.25 inch and width that varied between 2.5 to 3.0 inches for different specimens. The specimens were fatigue pre-cracked to a crack length, a, of 1.2 to 1.5 inch (a/W=0.45 to 0.5). The tests on plant trial material, discussed below, which did satisfy the validity criterion in ASTM E399 for K Ic , were conducted using compact tension specimens with a thickness, B=2.0 inch and width, W=4.0 inch. These specimens were fatigue pre-cracked to a crack length of 2.0 inch (a/W=0.5). All cases of comparison of data between alloys were made using results from specimens of the same size and under similar test conditions.  
       Example 1  
     Plant Trial—Plate  
       [0063]    A plant trial of the alloy of the present invention, designated alloy C80A, was conducted using a standard, full-size ingot cast with the composition (in wt. %): 7.35 Zn, 1.46 Mg, 1.64 Cu, 0.04 Fe, 0.02 Si and 0.11 Zr. The ingot was scalped, homogenized at 885° to 890° F. for 24 hours, and hot rolled to 6-inch thick plate. The rolled plate was solution heat treated at 885° to 890° F. for 140 minutes, spray quenched to ambient temperature, and cold stretched 1.5% for residual stress relieval. Sections from the plate were given a two-step aging comprising a 6-hour 250° F. first step aging followed by second step aging for 6, 8 and 11 hours, respectively, at 320° F. designated as times (t1), (t2) and (t3), respectively. Results from the tensile, fracture toughness, alternate immersion SCC, EXCO and electrical conductivity tests are presented in Table 2 below. FIG. 7 shows the cross plot of L-T plane-strain fracture toughness (K Ic ) versus longitudinal tensile yield strength TYS(L). A linear strength-toughness correlation trend was possible to define through the use of data from three different second-step aging times. Included in FIG. 7 are typical properties from 6-inch thick 7050-T7451 plates produced by industry specification, BMS 7-323C, and the 7040-T7451 typical values for 6-inch thick plates per AMS D99AA (Draft) specification (ref.  Preliminary Materials Properties Handbook ). The C80A alloy plate of the present invention clearly displays a much superior strength-toughness combination compared to either prior art 7050 or 7040 alloy plate. For example, with respect to 7050-T7451 plate, a TYS increase from about 64 ksi to 72 ksi (11% increase) is obtained at the equivalent K Ic  of about 35 ksi= 29  in. Alternatively, significant increase in KIc values is obtained at equivalent TYS. For example, with respect to 7040-T7451 a K Ic  (L−T) toughness increase from 32.3 to about 41 ksi{square root}in (28% increase) is obtained at the equivalent TYS (L) of 66.6 ksi.  
                                                                                           TABLE 2                           Properties of Plant Processed 6-inch Thick C80A Plates                                            SCC - Stress       Aging Time   L-UTS   L-TYS                       (ASTM G44)       at 320° F.   (T/4)   (T/4)   EL   L-CYS   L-TK Ic     EXCO   EC(T/4)   (20d-Pass)       (Hrs.)   (ksi)   (ksi)   (%)   (ksi)   (ksi{square root}in)   (T/4)   (% IACS)   (ksi)                    6 (t1)   77.1   74.9   6.8   73.2   33.6   EB   40.5   35       8 (t2)   75.6   72.5   7.3   71.0   35.2   EB   41.3   40       11 (t3)    71.9   67.2   8.6   65.6   40.5   EA   42.7   45                  
 
       Example 2  
     Plant Trial—Forging  
       [0064]    Die forging evaluation of the alloy of the present invention was performed in a plant-trial using two full-size production sheet ingots, designated C8XA-1 and C8XA-2, with the following compositions: Ingot C8XA-1: 7.35 Zn, 1.46 Mg, 1.64 Cu, 0.038 Fe, 0.022 Si and 0.11 Zr; Ingot C8XA-2: 7.39 Zn, 1.48 Mg, 1.91 Cu, 0.036 Fe, 0.024 Si and 0.11 Zr. A standard 7050 ingot was also run as a control. The ingots were homogenized at 885° F. for 24 hours and sawed to billets for forging. A closed die forged part was produced for evaluation of properties at three different thicknesses, namely, 2 inch, 3 inch and 7 inch thicknesses. The fabrication steps involved two pre-forming operations utilizing hand forging. This was followed by a blocker die operation and a final finish die operation both carried out in a 35,000 ton press. Forging temperatures used were about 725° to 750° F. The forged pieces were solution heat treated at 880° to 890° F. for 6 hours, quenched and cold worked 1 to 5% for residual stress relief. The parts were next given a 3-step aging treatment to develop a “T74” equivalent temper. The 7050 control pieces were aged to “T74” temper using the standard aging of 225° F. for 8 hours followed by 250° F. for 8 hours followed by 350° F. for 8 hours. Mechanical property tests were performed on samples taken from the same locations as used for the 7050 parts using standard procedures. Results from the tensile tests performed in longitudinal, long-transverse and short-transverse directions are presented in FIG. 8. In all three orientations, the tensile yield strength (TYS) for the alloy of the invention remained virtually unchanged for thicknesses ranging from 2 to 7 inches. In contrast, the properties for 7050 showed a drop in TYS values as the thickness increased from 2 to 3 to 7 inches consistent with the required minimum properties for this alloy (also shown in FIG. 8). Thus, the results in FIG. 8 demonstrate the advantages of the low quench sensitivity properties of the alloy forgings of the present invention in being insensitive to strength changes over a large thickness range, in contrast to the properties of the prior 7050 alloy.  
         [0065]    Hence, the present invention runs counter to conventional 7XXX series aluminum alloy design which indicates that a high Mg content is desirable for high strength. This is because a high Mg content is good for thin sections but not in thick product. It has been shown that more solute increases strength, but also concurrently increases the quench sensitivity of the material. The present invention recognizes that by lowering the Mg content to at or less than 1.7 wt. % and by controlling the Cu and Mg relationship to Mg≦Cu+0.3 or Mg≦Cu+0.3 or Mg≦Cu or Mg≦Cu (contrary to the teachings of U.S. Pat. No. 6,027,582), and by controlling the Zn content, the quench sensitivity of the alloy is decreased by keeping the Zn, Mg and Cu in solution at the slow cooling rate inherently present in thick gauges, viz., 4 to 8 inches or more. Thus, the alloy of the current invention seemingly runs counter to the historical alloy precedent in that, instead of increasing solute levels, the overall solute levels of Mg and Cu in the invention are actually decreased compared to current thick section alloys.  
         [0066]    Although the primary focus of the present invention is on product with thick cross sections quenched as rapidly as practical, those skilled in the art will recognize and appreciate that another application of the invention would be use of an intentionally slowed quench rate in parts with thin sections in order to reduce quench induced residual stresses, and distortion brought on by rapid quenching. Other, more quench sensitive alloys would suffer unacceptable losses in strength, toughness and other properties when treated this way (slowed quench rate) and would, thus, be unsuitable.  
         [0067]    Another variant of potential application arising from lower quench sensitivity of the present alloy is with parts that have both thick and thin sections. These parts will suffer less from differences in yield strength between the thick and thin cross sections caused by the quench rate differences. This would, in turn, reduce the incidence of bowing or distortion in the parts after stretching.  
         [0068]    It is known that, for a given 7XXX series alloy, as further artificial aging is progressively applied to the peak strength T6-type temper condition, the strength of the alloy progressively and systematically decreases while its fracture toughness and corrosion resistance progressively and systematically increase. Hence, those skilled in the art can select a specific temper condition with an optimum combination of strength, fracture toughness and corrosion resistance for a specific application. Indeed, such is the case for the alloy of the invention, as demonstrated in the cross plot of L-T plane strain fracture toughness K Ic  and L-yield strength, in FIG. 7, both measured at quarter plane (T/4) in the longitudinal direction for a 6-inch thick plate product. In FIG. 7, the alloy of the invention is shown to provide the combination of about 75 ksi yield strength with about 33 ksi{square root}in fracture toughness (at aging time t1 from Table 2), or about 72 ksi yield strength with about 38 ksi{square root}in fracture toughness (at aging time t2), or about 67 ksi yield strength with about 40 ksi{square root}in fracture toughness (at aging time t3). It is further understood by those skilled in the art that for a specific 7XXX series alloy, the strength-fracture toughness trend line can be interpolated and, to some extent, extrapolated to combinations of strength and fracture toughness beyond the three examples of alloy C80A given above and plotted in FIG. 7. The desired combination then can be accomplished by an appropriately selected artificial aging treatment.  
         [0069]    While the invention has been described largely in connection with aerospace structural applications, such is not necessarily intended to limit the invention which, it is believed, may find applications as a relatively thick cast, plate, extruded or forged product requiring a relatively high strength in a slowly quenched condition from the solution heat treating temperature. An example of one such application is a mold plate which is extensively machined into molds of various shapes to facilitate shaping or contouring processes in manufacturing processes. For this application, the desired material characteristics are high strength and low machining distortion. For alloys used as mold plates, a slow quench after solution heat treatment would be necessary to impart a low residual stress, which causes machining distortion. However, a slow quench also results in lower strength and other properties for existing 7XXX series alloys as a result of their higher quench sensitivity as previously described. It is the unique lack of quench sensitivity of the present alloy which permits a slow quench following solution heat treatment while still retaining the capability to attain a relatively high strength that makes it also an attractive choice for such non-structural applications as a thick mold plate.  
         [0070]    Having thus described the presently preferred embodiments of our invention, it will be understood by those skilled in the art that the invention may be otherwise embodied or modified without departing from the spirit and scope of the following claims.