The present invention is related to an aluminum alloy, and in particular, to an aluminum casting alloy that exhibits high strength and adequate ductility.
The use of aluminum alloys for lightweight components when compared to steel is known. For example, aluminum alloys are used extensively in the automotive and aircraft industry. In addition, with a face centered cubic structure, aluminum alloys have also found use in cryogenic type environments. For example, U.S. Pat. No. 7,060,139 filed on Nov. 8, 2002, titled “High Strength Aluminum Alloy Composition”, provides a high strength 7XXX series alloy that can be used at cryogenic temperatures and afford a tensile strength of at least 790 megapascals (MPa) with an elongation of at least 6 percent. Such an alloy can be used as part of a cryogenic pump, for example as an impeller, an inducer, and the like for a pump used to handle liquid nitrogen and/or liquid helium.
Although such an alloy can meet the demands of cryogenic type services, the alloy is a wrought material that requires one or more hot working steps in order to obtain a final component. In the alternative, if a cast aluminum alloy having sufficient strength and ductility could be provided, such cryogenic components could be cast without the need for additional hot working steps and thereby provide a more cost-effective material. However, high strength 7XXX series alloys are known to be difficult to cast, with most common defects resulting from the casting thereof being intergranular porosity, hot tears, and cold cracks. It is appreciated that the hot tears and subsequent cold cracks can result from tears in the solidification mushy zone resulting from the interplay between deformation of the partially coherent solid and a lack of interdendritic liquid feeding [1, 2]. As such, a foundry version of the high strength 7XXX series alloy that can be cast and provide suitable combination of strength and ductility without additional working would be desirable.
The present invention discloses a high strength cast Al—Zn—Mg—Cu (700 series) alloy, the cast alloy having a tensile strength of at least 500 MPa and 4% elongation at room temperature. In some instances, the cast aluminum alloy has a composition of about 5.5-9.0 weight percent (wt. %) of zinc, 2.0-3.5 wt. % of magnesium, 0.1-0.5 wt. % scandium, 0.05-0.20 wt % zirconium, 0.5-3.0 wt. % copper, 0.10-0.45 wt. % manganese, 0.01-0.35 wt. % iron, 0.01-0.20 wt. % silicon with a balance of aluminum and possible casting impurities. In other instances, the cast alloy has a tensile strength of at least 500 MPa and elongation at least 6% at room temperature.
A process for making the high strength aluminum cast alloy includes mixing and melting of aluminum and/or melting of an aluminum master alloy and providing alloying additions of zinc, magnesium, scandium, zirconium, copper, manganese, iron and/or silicon such that the above-described composition is provided. In addition, a melt of the alloy can be poured into a desired shape using any casting process known to those skilled in the art, illustratively including sand casting, investment casting, lost wax casting, and the like. The pouring temperature of the molten alloy should not be below 740° C. and a degassing step should be applied prior to pouring to remove dissolved gases (e.g. hydrogen) and solid inclusion from the melt. After the liquid alloy has been poured into the desired shape, it may or may not be allowed to solidify under atmospheric pressure. In some instances, the poured casting can be allowed to solidify under increased pressure, for example and for illustrative purposes only, 2 atmosphere, 5 atmosphere, 10 atmosphere, and the like. In addition, a cast component made from the cast aluminum alloy disclosed herein can be subjected to one or more heat treatment steps that afford improved mechanical properties thereof. For example, a heat treatment or temper for 7XXX series alloys known to those skilled in the art such as T1, T4, T6, T7 and the like can be performed on the cast component and provide improved mechanical properties when compared to a non-heat treated component. In addition, or in combination, heat treatment(s) that afford precipitation of coherent nano-particles of a Al3(Sc,Zr) phase can be performed to additionally improve mechanical properties.
The present invention discloses a high strength cast aluminum alloy. As such, the present invention has utility as a material of construction.
The high strength cast aluminum alloy can have a range of compositions that provide for a tensile strength of at least 500 megapascals (MPa) and an elongation of at least 2% at room temperature (Is this accurate? Is this what we want to claim?). In some instances, the cast aluminum alloy has a composition of about 5.5-9.0 weight percent (wt. %) zinc, 2.0-3.5 wt. % magnesium, 0.1-0.5 wt. % scandium, 0.05-0.20 wt. % zirconium, 0.5-3.0 wt. % copper, 0.10-0.45 wt. % manganese, 0.01-0.35 wt. % iron, 0.01-0.20 wt. % silicon and a balance of aluminum and possible casting impurities known to those skilled in the art.
A process for making the high strength cast aluminum alloy includes melting of aluminum and/or melting of an aluminum master alloy with alloying additions of zinc, magnesium, scandium, zirconium, copper, manganese, iron and/or silicon such that the above-described composition is provided. It is appreciated that the alloying additions can be added to a crucible containing the aluminum and/or aluminum master alloy before and/or after melting has been initiated.
Once a desired chemistry of the cast aluminum alloy has been obtained in the crucible, the liquid alloy can be degassed to remove hydrogen and poured into a desired shape using any casting process known to those skilled in the art, illustratively including sand casting, investment casting, lost wax casting, and the like. The pouring temperature of the molten alloy should not be less than 740° C. and generally should be in the range of 740-770° C. After pouring, the casting may or may not be solidified under atmospheric pressure (1 atm). Stated differently, the poured casting can be solidified at 1 atm, or in the alternative, solidified under 2 atm, 5 atm, 10 atm, and the like (1 atm=0.1013 MPa).
The solidified cast component can be subjected to a heat treatment and/or temper in order to precipitate nano-particles and improve the mechanical properties thereof. For example, heat treatments for 7XXX (series alloys as known to those skilled in the art, illustratively including T1, T4, T6, T7 and the like, can be performed on the cast component. In addition, or in combination, heat treatment(s) that afford precipitation of coherent nano-particles of a Al3(Sc,Zr) phase can also be performed to further improve mechanical properties.
In order to better illustrate the high strength cast aluminum alloy composition and a process for making a component out of the cast aluminum alloy, an example of the inventive alloy is provided below.
A cast aluminum (Al) alloy having the chemical composition shown in Table 1 was melted and cast at Universal Alloy Corporation, located in Anaheim, Calif. by direct chill (DC) casting in the form of 178 millimeter (mm) diameter billets.
After casting, the billets were homogenized by holding at 475° C. for 23 hours, then were slow cooled from 475 C to 250° C. within 14 hours and then cooled to room temperature in air. The homogenized billets were ultrasonically tested using an AMS standard AMS-STD-2154 N/C, Type I, Class A, and no cracks or porosity were detected. Tensile samples were extracted from the homogenized billets in three orthogonal directions. The longitudinal direction was parallel to the main axis of the billet. The radial direction was directed along the billet radius, and the transverse direction was perpendicular to the longitudinal and radial directions and the transverse samples were extracted from the billet section located half way from the billet surface to the billet center. These samples were heat treated to different tempers, shown in Table 2, and tensile tests were conducted at room and cryogenic temperatures. The results of the tensile tests are shown in Table 3 and Table 4. At room temperature (T=25° C.), the yield strength (YS) was in the range of 497-538 MPa, tensile strength was in the range of 586-605 MPa, and elongation (El) was above 6%. At the cryogenic temperature (T=−196° C.), YS=611-653 MPa, UTS=679-700 MPa, and El=2.0-3.1%.
Rectangular plates of 125 mm long, 32 mm wide and with the thicknesses of 6.3 mm (Plate #1) and 12.7 mm (Plate #2) were produced by sand casting at Eck Industries. The direct chill casting billets described in Example 1 were used as melt stock for this sand casting with pieces of the billet re-melted, the molten alloy degassed and the poured into sand molds. The pouring temperature was 760° C. The mold to produce the 12.7 mm thick plate had a steel chill plate at the bottom of the mold for faster solidification. After solidification, the plates had equiaxed dendritic structures with the grain size of 153 μm for plate #1 and 33 μm for plate #2. Several plates were conventionally HIP'd at Kittyhawk Products. The HIP cycle consisted of heating to 521° C. with a continuous increase in the pure argon pressure, holding at this temperature for 2 hours at the pressure of 104 MPa, and slow cooling to room temperature. Kittyhawk Products use this HIP cycle in their practice to close shrinkage pores in Al—Si alloy castings.
The plates were solution treated at 480° C. for 8 hours, water quenched and aged at 120° C. for 20 hours (T6 temper). Tensile specimens were extracted from the plates and tested at room temperature (RT) in accord to ASTM B557-10 standard. The results are given in Table 5 and Table 6. The YS was above 500 MPa and UTS was ˜550 MPa for cast and T6 tempered plates and ˜590-600 MPa for HIPd and T6 tempered plates. Using HIP increased tensile ductility (elongation) of the plates considerably, from ˜1.5-2.2% to 11-12%.
In this example, charges of the DC cast billets described in Example 1, of approximately 11.5 kilograms (kg) each, were re-melted in an electrically heated crucible furnace at 760° C., degassed using argon gas for 3 minutes, and poured into two wedge sand molds. One mold was kept in air under the normal atmospheric pressure, while another mold was immediately placed inside a pressure vessel. The pressure vessel was 1.2 meter (m) in diameter and 1.8 m in height, and pressure was applied via a near-equal mixture of dry compressed air and compressed nitrogen gas. The peak pressure of 10 atm was achieved within the pressure vessel approximately 90 seconds after pouring of a wedge. In this manner, one cast wedge of the Al alloy was solidified at 1 atm and another wedge was solidified at 10 atm of pressure.
A schematic of the wedge casting and gating system is shown in
As known to those skilled in the art, variation in wedge thickness produces a variation in solidification rate with the fastest cooling rate occurring at the bottom of the wedge (thinnest section) and decreasing cooling rates occurring at increasing thicknesses. To determine actual cooling rates, thermocouples were placed at four different wedge thickness locations designated A, B, C and D. All four of the thermocouples were placed at the center of the wedge and their location relative to the height of the wedge and the section thickness for a particular height are given in Table 7.
Microstructures of the sand cast alloy samples were studied using scanning electron microscopy (SEM) and analyzed using Fovea Pro image processing software. In addition, subsequent hot isostatic pressing (HIP) of a wedge casting solidified at 10 atm was conducted at Bodycote North American HIP located in Princeton, Ky. The HIP processing employed temperature ramping of 4° C./min up to 460° C. with a simultaneous increase in pressure up to 207 MPa, followed by holding at 460° C. and 207 MPa for 2 hours, ramping to 500° C. while maintaining the pressure at 207 MPa, holding for 1 hour, and furnace cooling at release pressure.
Tensile samples were extracted from different sections of as-cast wedges solidified under 1 atm, 10 atm and 10 atm plus HIPing, and the tensile tests were conducted at room temperature in accordance to ASTM standards ASTM E8-04 and ASTM B557-10, independently at the University of Alabama at Birmingham (UAB) and at the Air Force Research Laboratory (AFRL) using servo-hydraulic materials testing systems (MTS 810). The tensile samples tested at UAB had a cylindrical gauge shape with a gauge length of 36 mm and a gauge diameter of 9 mm. In addition, the UAB tensile samples were tested at a constant ramp speed of 0.036 mm/sec (initial strain rate of 0.001 s−1). Tensile samples tested at AFRL had a rectangular cross-section gauge section of 2.5×3.6 mm, a gauge length of 20 mm and were tested at a constant ramp speed of 0.02 mm/sec (initial strain rate of 0.001 s−1). Prior to testing, the samples were heat treated to T6, T7 or T4 tempers. The solution treatment and aging conditions for the tempers are shown in Table 8 below and the solution treated samples were water quenched prior to aging.
Turning now to
In particular, within the first 400 seconds, solidification occurred fastest in the thinnest section A, followed by slower solidification in section B, and then the slowest solidification in sections C and D which exhibited essentially identical solidification kinetics. After 400 seconds from pouring, the temperatures in sections A, B, C and D were 532° C., 571° C., 601° C. and 610° C., respectively, for a 1 atm solidified casting and 522° C., 568° C., 600° C., and 600° C., respectively, for a 10 atm solidified casting. It is interesting to note that at longer times the cooling rate slowed down in thinner sections, accelerated in thicker sections and differences in temperature between the four sections continuously decreased. In fact, after approximately 1000 seconds of cooling, all of the sections cooled at approximately 0.1-0.15° C./sec.
The temperature at which solidification of the alloy started (i.e. the liquidus temperature, TL) was clearly identified at the end of a rapid decrease in cooling rate near the beginning of the cooling rate versus temperature curve. In addition, the liquidus temperature was almost insensitive to cooling rate and applied pressure (TL=633±2° C.) as shown by the TL values for sections A-D and solidification pressures of 1 atm and 10 atm in Table 9.
Solidification of the alloy ended with formation of a eutectic which exhibited a rapid but temporary decrease in cooling rate for sections B, C and D. The solidus temperature (TS) of the alloy in sections B-D are given in Table 4 and the average value was determined to be TS469±2° C. In contrast, the cooling curve from the thinnest section A did not indicate a eutectic reaction, i.e. the cooling rate did not show a rapid drop below 550° C. Not being bound by theory, the rapid solidification of section A at temperatures above 550° C. could have led to less eutectic-forming elements remaining in the liquid as compared to sections B-D and as such a considerably reduced volume fraction of eutectic solidification in section A.
Macrographs of the transverse cross-sections of sand cast wedge samples solidified under 1 atm and 10 atm are shown in
Micrographs illustrating the microstructure of the four sections A-D for as-cast wedges solidified at 1 atm and 10 atm are shown in
In contrast to grain size, volume fraction of second phase particles did depend on solidification pressure and also depended on sample thickness. For example, for the wedge solidified at 1 atm, the volume fraction of second phase particles was approximately 6.8% in Section A (12.7 mm) and continuously decreased to about 2.1% in section D (50.8 mm). However, for the wedge soldified at 10 atm, the volume fraction of second phase particles was approximately the same, 3.5% in Sections A, B, and C and slightly decreased to about 2.8% in Section D (see
Regarding volume fraction of pores, for wedges solidified at 1 atm, pores increased from about 1.7% to 4% with an increase in wedge thickness from 12 mm to 51 mm (
In addition to determining the effect of pressure on solidification of the inventive cast alloy, the effect of HIP processing was studied by taking an 80 mm wide edge piece of the 10 atm cast wedge and subjecting it to HIP parameters/conditions described above. Microstructures of HIPd samples were analyzed in wedge thickness regions which approximately corresponded to the thicknesses of Sections A-D. SEM images of the microstructures are shown in
Although not readily apparent from comparing
In contrast, HIPing of the material did decrease the volume fraction of intergranular second phase particles as shown in
The positive effect of solidification pressure on mechanical properties for the cast aluminum alloy is shown by comparing the room temperature tensile properties in Table 10 and Table 11. As shown in the tables, the tensile samples were subjected to T6 and T7 tempers before testing. In 13 to 36 mm thick sections solidified at 10 atm, the yield strength (YS) was 23-28% greater and the ultimate tensile strength (UTS) was 21-33% greater than corresponding sections solidified at 1 atm. In addition, for wedge sections generally 33 mm thick and thinner, the 10 atm cast alloy exhibited a YS of between 489-522 MPa and a UTS of between 529-592 MPa. It is appreciated that heretofor cast aluminum alloys have not provided such high strength values [3].
Tensile ductility of the 1 atm cast alloy was very low, about 0.8-1.5%, and was practically unaffected by the wedge thickness. The low ductility is associated with gas porosity present in this sample throughout the thickness. In contrast, solidification at 10 atm led to a noticeable increase in tensile ductility for wedge thicknesses up to 30 mm where gas porosity was not developed. In thicker regions, the ductility rapidly decreased due to gas porosity present in these thicker sections.
The 10 atm cast alloy also exhibited good response to natural aging. For example, after holding at room temperature for 7 days, 10 to 20 mm thick sections had yield strengths of between 350-360 MPa, ultimate tensile strengths of between 520-540 MPa and elongation values above 13% (see Table 12). The mechanical properties rapidly decreased to YS=281 MPa, UTS=380 MPa and El=5.8% with an increase in the wedge thickness to 37 mm; however, the T4-tempered samples were much more ductile than T6-tempered samples in the studied thickness range. It is appreciated that finer GP-I zones can form during natural aging (T4 temper) and thereby result in a more homogeneous deformation of grains and reduced stress concentrations at grain boundaries when compared to the presence of coarser GP-II zones and η′ particles formed during artificial aging (T6 and T7 tempers).
Brinell hardness of the cast alloy increased with an increase in solidification pressure and a decrease in the wedge thickness (compare Table 10 and Table 12). After T6 and T7 tempers, 13 to 25 mm thick sections of the 1 atm cast alloy had hardness between 138-145 BHN, compared to a hardness of between 152-165 BHN for the 10 atm cast alloy. The T4 tempered samples had hardness values of between 114-125 BHN for 13-36 mm thick wedge sections.
As shown in Table 13 below, the YS slightly decreased, whereas UTS, ductility and hardness increased for the 10 atm cast plus HIPd plus T6 or T7 tempered 13-30 mm thick wedge samples when compared to the 10 atm cast plus T6 or T7 tempered wedge samples (no HIP) of the same thickness (compare Table 13 with Table 11). The slight decrease in YS could be due to coarsening of Al3(Sc,Zr) nanoparticles during HIPing, while the increase in the UTS could be associated with strain hardening and increased ductility of the HIPd samples.
Table 15 shows typical mechanical properties of these cast alloys at room temperature. The data is taken from ref. [3].
It is appreciated that currently available Al—Cu—Mg alloys have the highest strength capabilities among all commercial Al alloy castings [3]. In particular, the 201.0-T7 alloy has a UTS=470 MPa, YS=415 MPa and El=6%. However, these alloys are also susceptible to solidification cracking and interdendritic shrinkage. Exacting foundry techniques are required to avoid these conditions. In addition, the Al—Cu based alloy castings are susceptible to general corrosion and stress corrosion.
Al—Si based alloy castings are the most widely used sand cast alloys. They exhibit good fluidity, castability and corrosion resistance and such alloys containing Cu and Mg are heat treatable. However, tensile strength and yield strength of these alloys are relatively low. For example, maximum strength values of UTS=345 MPa and YS=290 MPa are achieved in the 359.0 alloy. The tensile ductility of these alloys varies from <1% to 6%, depending on the composition and heat treatment [3].
The castability of currently available Al—Zn—Mg based alloys is poor and good foundry practices are required to minimize hot tearing and shrinkage defects. These alloys typically display moderate tensile properties. For example, the 771.0-T7 alloy casting exhibits maximum strength levels among the Al—Zn—Mg based alloy castings compositions [3], namely a UTS=330 MPa and a YS=310 MPa. In addition, hardness is in the range of 60-135 HB and elongation is in the range of 1-5%, depending on the composition and heat treatment (see Table 15). The Al—Zn—Mg based castings also show good machinability and corrosion resistance [3].
The proposed aluminum alloy casting composition is an Al—Zn—Mg based casting alloy. It is different from the known casting compositions by the presence of Sc and Zr, as well as by different combinations of other elements.
The room temperature tensile strength of the proposed alloy casting composition is much higher than the strength of any commercially available aluminum alloy castings. In particular, the tensile strength of the new alloy composition is at least 75% higher than that of the 771.0 alloy and at least 17% higher than the 201.0.
Fluidity of the proposed aluminum alloy casting composition (SSA008) is similar to the fluidity of A356.0, which is one of the best tastable aluminum alloys (Table 16). The fluidity indices of the molten alloys were determined using a N-Tec fluidity mold and three superheat temperatures, ΔT=T−TL, of 20° C., 60° C. and 110° C., for both A356.0 and SSA008 alloys. Here T is the temperature of the molten alloy during pouring and TL is the liquidus temperature, TL=612° C. for A356.0 and TL=633° C. for SSA008. The fluidity mold had five fingers of the same length and different thicknesses, and the mold was preheated on a hot plate to approximately 300° C. before liquid alloy was poured therein. The fluidity index at each superheat temperature was determined as the total length of the solidified alloy in the mold fingers with results from the fluidity tests provided in Table 16. As show by the data, the developmental alloy and the commercial A356.0 alloy exhibit similar fluidity.
In view of the teaching presented herein, it is to be understood that numerous modifications and variations of the present invention will be readily apparent to those of skill in the art. For example, while the invention has primarily been described with specific alloying additions to be made to aluminum, it is appreciated that other alloying additions known to those skilled in the art can be included and fall within the scope of the invention. In addition, casting impurities that occur during the casting of aluminum alloys can be present within a cast component. As such, the foregoing is illustrative of specific embodiments of the invention, but is not meant to be a limitation upon the practice thereof. Therefore, the specification should be interpreted broadly.
This application claims priority to U.S. Provisional Patent Application No. 61/392,310 filed on Oct. 12, 2010, having the same title and which is incorporated herein in its entirety by reference.
Number | Date | Country | |
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61392310 | Oct 2010 | US |