Patent Grant
8002913
This invention relates to an AA7000-series alloy comprising 3 to 10% Zn, 1 to 3% Mg, at most 2.5% Cu, Fe <0.25%, and Si >0.12 to 0.35%, and to a method of manufacturing these aluminum alloy products. More particularly, the invention relates to aluminum wrought products in relatively thick gauges, in particular i.e. about 30 to 300 mm thick. While typically practiced on rolled plate product forms, this invention may also find use with manufacturing extrusions or forged product shapes. Representative structural component parts made from the alloy product include integral spar members and the like which are machined from thick wrought sections, including rolled plate. This invention is particularly suitable for manufacturing high strength extrusions and forged aircraft components. Such aircraft include commercial passenger jetliners, cargo planes and certain military planes. In addition, non-aerospace parts like various thick mould plates or tooling plates may be made according to this invention.
As will be appreciated herein below, except as otherwise indicated, alloy designations and temper designations refer to the Aluminum Association designations in Aluminum Standards and Data and the Registration Records, as published by the Aluminum Association in 2006.
For any description of alloy compositions or preferred alloy compositions, all references to percentages are by weight percent unless otherwise indicated.
Different types of aluminum alloys have been used in the past for forming a variety of products for structural applications in the aerospace industry. Designers and manufacturers in the aerospace industry are constantly trying to improve fuel efficiency, product performance and constantly trying to reduce the manufacturing and service costs. The preferred method for achieving the improvements, together with the cost reduction, is the uni-alloy concept, i.e. one aluminum alloy that is capable of having improved property balance in the relevant product forms.
State of the art at this moment is high damage tolerant AA2x24 (i.e. AA2524) or AA6x13 or AA7x75 for fuselage sheet, AA2324 or AA7x75 for lower wing, AA7055 or AA7449 for upper wing and AA7050 or AA7010 or AA7040 or AA7140 for wing spars and ribs or other sections machined from thick plate. The main reason for using different alloys for each different application is the difference in the property balance for optimum performance of the whole structural part.
For fuselage skin, damage tolerant properties under tensile loading are considered to be very important, that is a combination of fatigue crack growth rate (“FCGR”), plane stress fracture toughness and corrosion. Based on these property requirements, high damage tolerant AA2×24-T351 (see e.g. U.S. Pat. No. 5,213,639 or EP-1026270-A1) or Cu containing AA6xxx-T6 (see e.g. U.S. Pat. No. 4,589,932, U.S. Pat. No. 5,888,320, US-2002/0039664-A1 or EP-1143027-A1) would be the preferred choice of civilian aircraft manufactures.
For lower wing skin a similar property balance is desired, but some toughness is allowably sacrificed for higher tensile strength. For this reason AA2x24 in the T39 or a T8x temper are considered to be logical choices (see e.g. U.S. Pat. No. 5,865,914, U.S. Pat. No. 5,593,516 or EP-1114877-A1).
For upper wing, where compressive loading is more important than the tensile loading, the compressive strength, fatigue (SN-fatigue or life-time or FCGR) and fracture toughness are the most critical properties. Currently, the preferred choice would be AA7150, AA7055, AA7449 or AA7x75 (see e.g. U.S. Pat. No. 5,221,377, U.S. Pat. No. 5,865,911, U.S. Pat. No. 5,560,789 or U.S. Pat. No. 5,312,498). These alloys have high compressive yield strength with at the moment acceptable corrosion resistance and fracture toughness, although aircraft designers would welcome improvements on these property combinations.
For thick sections having a thickness of more than 3 inch or parts machined from such thick sections, a uniform and reliable property balance through thickness is important. Currently, AA7050 or AA7010 or AA7040 (see U.S. Pat. No. 6,027,582) or AA7085 (see e.g. US Patent Application Publication No. 2002/0121319-A1 and U.S. Pat. No. 6,972,110) are used for these types of applications. Reduced quench sensitivity, that is deterioration of properties through thickness with lower quenching speed or thicker products, is a major wish from the aircraft manufactures. Especially the properties in the ST-direction are a major concern of the designers and manufactures of structural parts.
A better performance of the aircraft, i.e. reduced manufacturing cost and reduced operation cost, can be achieved by improving the property balance of the aluminum alloys used in the structural part and preferably using only one type of alloy to reduce the cost of the alloy and to reduce the cost in the recycling of aluminum scrap and waste.
Accordingly, it is believed that there is a demand for an aluminum alloy capable of achieving the improved proper property balance in almost every relevant product form.
It is an object of the present invention to provide AA7000-series alloys having improved property balance.
It is another object of the present invention to provide a wrought aluminum alloy product of an AA7000-series alloy comprising 3 to 10% Zn, 1 to 3% Mg, at most 2.5% Cu, Fe <0.25%, and Si >0.12 to 0.35% having improved properties, in particular having improved fracture toughness.
It is another object of the present invention to provide a method of manufacturing such improved AA7000-series alloy products.
These and other objects and further advantages are met or exceeded by the present invention method of manufacturing a wrought aluminum alloy product of an AA7000-series alloy comprising Si >0.12 to 0.35%, and preferably comprising 3 to 10% Zn, 1 to 3% Mg, at most 2.5% Cu, Fe <0.25%, and Si >0.12 to 0.35%, the method comprising the steps of:
According to this invention there is at least one heat treatment carried out at a temperature in a range of more than 500° C. but lower than the solidus temperature of the subject AA7000 aluminum alloy, and wherein this heat treatment is carried out either: (i) after the homogenisation heat treatment but prior to hot working, or (ii) after the solution heat treatment of step e.), or (iii) both after the homogenisation heat treatment but prior to hot working and also after the solution heat treatment of step e.).
The aluminum alloy can be provided as an ingot or slab or billet for fabrication into a suitable wrought product by casting techniques regular in the art for cast products, e.g. DC-casting, EMC-casting, EMS-casting. Slabs resulting from continuous casting, e.g. belt casters or roll casters, also may be used, which in particular may be advantageous when producing thinner gauge end products. Grain refiners such as those containing titanium and boron, or titanium and carbon, may also be used as is well-known in the art. After casting the alloy stock, the ingot is commonly scalped to remove segregation zones near the cast surface of the ingot.
It is known in the art that the purpose of a homogenisation heat treatment has the following objectives: (i) to dissolve as much as possible coarse soluble phases formed during solidification, and (ii) to reduce concentration gradients to facilitate the dissolution step. A preheat treatment achieves also some of these objectives. A typical preheat treatment for AA7000-series alloys would be a temperature of 420 to 460° C. with a soaking time in the range of 3 to 50 hours, more typically for 3 to 20 hours.
Firstly, the soluble eutectic phases such as the S-phase, T-phase, and M-phase in the alloy stock are dissolved using regular industry practice. This is typically carried out by heating the stock to a temperature of less than 500° C., and typically in a range of 450 to 485° C., as the S-phase eutectic phase (Al2MgCu-phase) has a melting temperature of about 489° C. in AA7000-series alloys and the M-phase (MgZn2-phase) has a melting point of about 478° C. As is known in the art this can be achieved by a homogenisation treatment in said temperature range and allowing the stock to cool to the hot working temperature, or after homogenisation the stock is subsequently cooled and reheated to hot working temperature. The regular homogenisation process can also be done in two or more steps if desired, and which are typically carried out in a temperature range of 430 to 490° C. for AA7000-series alloys. For example in a two step process, there is a first step between 457 and 463° C., and a second step between 470 and 485° C., to optimise the dissolving process of the various phases depending on the exact alloy composition.
The soaking time at the homogenisation temperature according to industry practice is alloy dependent as is well known to the skilled person, and is commonly in the range of about 1 to 50 hours. The heat-up rates that can be applied are those which are regular in the art.
This is where the homogenisation practice according to the prior art stops. However, it is an important aspect of the present invention that after the regular homogenisation practice where the alloy composition allows complete dissolution of soluble phases (eutectics) present from solidification at least one further heat treatment can be carried out at a temperature in a range of more than 500° C. but at a temperature lower than the solidus temperature of the subject alloy.
For the AA7000-series alloys the preferred temperature is in a range of >500 to 550° C., preferably 505 to 540° C., and more preferably 510 to 535° C., and more preferably of at least 520° C.
For the alloy system the soaking time at this further heat treatment is from about 1 to up about 50 hours. A more practical soaking time would not be more than about 30 hours, and preferably not more than about 15 hours. A too long soaking time may lead to an undesired coarsening of dispersoids adversely affecting the mechanical properties of the final alloy product.
The skilled person will immediately recognise that at least the following alternative homogenisation practices can be used, while achieving the same technical effect:
In the alternatives wherein following the heat treatment according to this invention the stock is firstly cooled to, for example, ambient temperature prior to reheating for hot working, preferably a fast cooling rate is used to prevent or at least minimise uncontrolled precipitation of various secondary phases, e.g. Al2CuMg or Al2Cu or Mg2Zn.
Following the preheat and/or homogenisation practice according to this invention the stock can be hot worked by one or more methods selected from the group consisting of rolling, extrusion, and forging, preferably using regular industry practice. The method of hot rolling is preferred for the present invention.
The hot working, and hot rolling in particular, may be performed to a final gauge, e.g. 3 mm or less or alternatively thick gauge products. Alternatively, the hot working step can be performed to provide stock at intermediate gauge, typical sheet or thin plate. Thereafter, this stock at intermediate gauge can be cold worked, e.g. by means of rolling, to a final gauge. Depending on the alloy composition and the amount of cold work an intermediate anneal may be used before or during the cold working operation.
In an embodiment of the method according to this invention following the regular practice of SHT and fast cooling for the subject aluminum alloy product, the stock is subjected to the further heat treatment according to this invention, one may designate this as a second SHT, at a higher temperature than the first regular SHT, where after the stock is rapidly cooled to avoid undesirable precipitation out of various phases. Between the first and second SHT the stock can be rapidly cooled according to regular practice, or alternatively the stock is ramped up in temperature from the first SHT to the second SHT and after a sufficient soaking time it is subsequently rapidly cooled. This second SHT is to further enhance the properties in the alloy products and is preferably carried out in the same temperature range and time range as the homogenisation treatment according to this invention as set out in this description, together with the preferred narrower ranges. However, it is believed that also shorter soaking times can still be very useful, for example in the range of about 2 to 180 minutes. This further heat treatment may dissolve as much as practically possible any of the Mg2Si phases which may have precipitated out during cooling from the homogenisation treatment or the during a hot working operation or any other intermediate thermal treatment. The solution heat treatment is typically carried out in a batch furnace, but can also be carried out in a continuous fashion. After solution heat treatment, it is important that the aluminum alloy be cooled to a temperature of 175° C. or lower, preferably to ambient temperature, to prevent or minimise the uncontrolled precipitation of secondary phases, e.g. Al2CuMg and Al2Cu, and/or Mg2Zn. On the other hand cooling rates should preferably not be too high in order to allow for a sufficient flatness and low level of residual stresses in the product. Suitable cooling rates can be achieved with the use of water, e.g. water immersion or water jets.
Yet, in a further embodiment of this invention the defined AA7000-series alloy products are processed using regular homogenisation and/or preheat practice, and where after the products are processed using the preferred SHT as set out above, thus regular SHT followed by the second solution heat treatment in the defined temperature and time range, together with the preferred narrower ranges. This will result in the same advantages in product properties. It is possible to carry out the first regular SHT followed by rapid cooling and reheating to the soaking temperature of the second SHT, alternatively the temperature is ramped up from the first to the second SHT and after a sufficient soaking time it is subsequently rapidly cooled.
The stock may be further cold worked, for example, by stretching in the range of about 0.5 to 8% of its original length to relieve residual stresses therein and to improve the flatness of the product. Preferably the stretching is in the range of about 0.5 to 6%, more preferably of about 0.5 to 5%.
After cooling the stock is aged, typically at ambient temperatures, and/or alternatively the stock can be artificially aged. The artificial ageing can be of particular use for higher gauge products. Depending on the alloy system this ageing can de done by natural ageing, typically at ambient temperatures, or alternatively by means of artificially ageing. All ageing practices known in the art and those which may be subsequently developed can be applied to the AA7000-series alloy products obtained by the method according to this invention to develop the required strength and other engineering properties.
A desired structural shape is then machined from these heat treated plate sections, more often generally after artificial ageing, for example, an integral wing spar. SHT, quench, optional stress relief operations and artificial ageing are also followed in the manufacture of thick sections made by extrusion and/or forged processing steps.
The effect of the heat treatment according to this invention is that the damage tolerance properties are improved of the alloy product compared to the same aluminum alloy having also high Si content but processed without this practice according to the present invention. In particular an improvement can be found in one or more of the following properties: the fracture toughness, the fracture toughness in S-L orientation, the fracture toughness in S-T orientation, the elongation at fracture, the elongation at fracture in ST orientation, the fatigue properties, in particular FCGR, S-N fatigue or axial fatigue, the corrosion resistance, in particular exfoliation corrosion resistance, or SCC or IGC. It has been shown that there is a significant enhancement in mechanical properties of as much as 15%, and in the best examples of more than 20%.
In addition, similar enhanced properties are achieved, or at least not adversely affected, with the aluminum alloy products according to this invention and preferably processed according to this invention compared to the same alloy composition but having the regular low Si content and processed according to regular industry practice. This would allow the manufacturing of aluminum alloy product having similar or equivalent properties compared to the low Si alloys, but in a more cost effective manner as source material having a low Si-content is more expensive.
The following explanation for the surprisingly improved properties of the wrought product of this invention is put forward, with the caveat that it is merely an expression on belief and does not presently have complete experimental support.
The prior art refers to the Mg2Si constituent phase as being insoluble in AA7000-series aluminum alloys and these particles are known fatigue initiation sites. In particular for aerospace applications, the prior art indicates that the Fe and Si content need to be controlled to very low levels to provide products with improved damage tolerant properties such as Fatigue Crack Growth Rate resistance (“FCGR”) and fracture toughness. From various prior art documents it is clear that the Si content is treated as an impurity and should be kept at a level as low as reasonably possible. For example US-2002/0121319-A1, incorporated herein by reference, discusses the impact of these impurities on the alloying additions and states that Si will tie up some Mg thereby leaving an “Effective Mg” content available for solution, it is suggested that this be remedied by additional additions of Mg to compensate for the Mg tied up with the Mg2Si, see section [0030] of US-2002/0121319-A1. However, at no point it is suggested that the Mg2Si could be reintroduced into solution by a controlled heat treatment practice. With regard to the homogenisation practice it is mentioned that homogenisation may be conducted in a number of controlled steps but ultimately state that a preferred combined total volume fraction of soluble and insoluble constituents be kept low, preferably below 1% volume, see section [0102] of US-2002/0121319-A1. Within the examples, times and temperatures of heat treatments are given but at no point are the temperatures or times disclosed adequate in attempting the dissolution of Mg2Si constituent particles, i.e. homogenisation temperature of up to 900° F. (482° C.) and solution treatment temperature of up to 900° F. (482° C.).
However, it has been found in accordance with the invention that for various AA7000-series aluminum alloys, the generally perceived constituent phase Mg2Si is soluble via carefully controlled heat treatment and if they cannot be taken in complete solution then their morphology can be spherodised in such a way that fatigue and/or fracture toughness properties are improved. Once in solid solution, most of the Si and/or Mg will be available for subsequent ageing that may further enhance mechanical and corrosion properties. By deliberately increasing the Si content in the alloys according to this invention more of this Si is available for subsequent ageing practices but without having the detrimental coarse Mg2Si phases in the final product. The gained improvements by the purposive addition of Si could also be sacrificed to some extent by making the alloy composition leaner in Mg and/or Cu thus improving the toughness of the alloy product. Thus the generally perceived detrimental impurity element Si is now being converted into a purposive alloying element having various advantageous technical effects.
For the AA7000-series alloys the upper limit for the Si content is about 0.35%, and preferably of about 0.25%, as a too high Si content may result in the formation of too coarse Mg2Si phases which cannot be taken in complete solid solution and thereby adversely affecting the property improvements gained. For the AA7000-series alloys the lower limit for the Si-content is >0.12%. For these alloy systems a more preferred lower limit for the Si-content is about 0.15%, and furthermore preferably about 0.17%.
A wrought AA7000-series alloy product that can be processed favorably according to the method of this invention, comprises, in wt. %:
Zn about 3 to 10%
Mg about 1 to 3%
Cu 0 to about 2.5%
Fe <0.25%, preferably <0.10%
Si >0.12 to 0.35%, preferably >0.12 to 0.25%, more preferably about 0.15 to 0.25%,
one or more elements selected from the group consisting of:
the alloy optionally containing at most:
In a preferred embodiment the alloys processed using the method according to this invention have a lower limit for the Zn-content of about 5.5% and preferably about 6.1%, and more preferably of about 6.4%. And a more preferred upper limit for the Zn content is about 8.5%, and more preferably about 8.0%.
In a preferred embodiment the alloys processed using the method according to this invention have a preferred upper limit for the Mg content of about 2.5%, and preferably about 2.0%, and more preferably of about 1.85%.
In a preferred embodiment the alloys processed using the method according to this invention have a lower limit for the Cu-content of about 0.9% and more preferably about 1.1%. A more preferred upper limit for the Cu content is about 2.1%, and more preferably about 1.9%.
Traditionally, beryllium additions have served as a deoxidizer/ingot cracking deterrent. Though for environmental, health and safety reasons, more preferred embodiments of this invention are substantially Be-free. Minor amounts of Ca and Sr alone or in combination can be added to the alloy for the same purposes as Be.
The Fe content for the alloy should be less than 0.25%. When the alloy product is used for aerospace application preferably the lower-end of this range is preferred, e.g. less than about 0.10%, and more preferably less than about 0.08% to maintain in particular the toughness at a sufficiently high level. Where the alloy product is used for tooling plate application, a higher Fe content can be tolerated. However, it is believed that also for aerospace application a moderate Fe content, for example about 0.09 to 0.13%, or even about 0.10 to 0.15%, can be used. Although the skilled person would believe that this has an adverse effect on the toughness of the product, some of this loss in properties, if not all, is gained back when using the method according to this invention. The resultant would be an alloy product, although having moderate Fe levels, but when processed according to this invention it has properties equivalent to the same alloy product except for a lower Fe content, e.g. 0.05 or 0.07%, when processed using regular practice. Thus similar properties are achieved at higher Fe-levels, which has a significant cost advantage as source material having very low Fe-contents is expensive.
Silver in a range of at most about 0.5% can be added to further enhance the strength during ageing. A preferred lower limit for the Ag addition would be about 0.03% and more preferably about 0.08%. A preferred upper limit is about 0.4%.
Each of the dispersoid forming elements Zr, Sc, Hf, V, Cr and Mn can be added to control the grain structure and the quench sensitivity. The optimum levels of dispersoid formers depend on the processing, but when one single chemistry of main elements (Zn, Cu and Mg) is chosen within the preferred window and that chemistry will be used for all relevant products forms, then Zr levels are less than about 0.5%.
A preferred maximum for the Zr level is 0.2%. A suitable range of the Zr level is about 0.03 to 0.20%. A more preferred upper-limit for the Zr addition is about 0.15%. Zr is a preferred alloying element in the alloy product when processed according to this invention. Although Zr can be added in combination with Mn, for thicker gauge products manufactured using the method of this invention it is preferred that when Zr is added that any addition of Mn is avoided, preferably by keeping Mn at a level of less than 0.03%. In thicker gauge product the Mn phases coarsens more rapid than the Zr phases, thereby adversely affecting the quench sensitivity of the alloy product.
The addition of Sc is preferably not more than about 0.5% or more preferably not more than 0.3%, and even more preferably not more than about 0.18%. When combined with Sc, the sum of Sc+Zr should be less then 0.3%, preferably less than 0.2%, and more preferably at a maximum of about 0.17%, in particular where the ratio of Zr and Sc is between 0.7 and 1.4%.
Another dispersoid former that can be added, alone or with other dispersoid formers is Cr. Cr levels should preferably be below about 0.4%, and more preferably at a maximum of about 0.3%, and even more preferably about 0.2%. A preferred lower limit for the Cr would be about 0.04%. Although Cr alone may not be as effective as solely Zr, at least for use in tooling plate of the alloy wrought product, similar hardness results may be obtained. When combined with Zr, the sum of Zr+Cr should not be above about 0.23%, and preferably not more than about 0.18%.
The preferred sum of Sc+Zr+Cr should not be above about 0.4%, and more preferably not more than 0.27%.
In another embodiment of the aluminum alloy wrought product according to the invention the alloy product is free of Cr, in practical terms this would mean that the Cr content is at regular impurity levels of <0.05%, and preferably <0.02%, and more preferably the alloy is essentially free or substantially free from Cr. With “substantially free” and “essentially free” we mean that no purposeful addition of this alloying element was made to the composition, but that due to impurities and/or leaching from contact with manufacturing equipment, trace quantities of this element may, nevertheless, find their way into the final alloy product. In particular for thicker gauge products (e.g. more than 3 mm) the Cr ties up some of the Mg to form Al12Mg2Cr particles which adversely affect quench sensitivity of the wrought alloy product, and may form coarse particles at the grain boundaries thereby adversely affecting the damage tolerance properties.
Mn can be added as a single dispersoid former or in combination with one of the other dispersoid formers. A maximum for the Mn addition is about 0.4%. A suitable range for the Mn addition is in the range of about 0.05 to 0.4%, and preferably in the range of about 0.05 to 0.3%. A preferred lower limit for the Mn addition is about 0.12%. When combined with Zr, the sum of Mn plus Zr should be less then about 0.4%, preferably less than about 0.32%, and a suitable minimum is about 0.12%.
In another embodiment of the aluminum alloy product according to the invention the alloy is free of Mn, in practical terms this would mean that the Mn-content is <0.03%, and preferably <0.02%, and more preferably the alloy is essentially free or substantially free from Mn. By “substantially free” and “essentially free” we mean no purposeful addition of this alloying element was made to the composition, but that due to impurities and/or leaching from contact with manufacturing equipment, trace quantities of this element may, nevertheless, find their way into the final alloy product.
In another preferred embodiment of the aluminum alloy wrought product according to this invention, the alloy has no deliberate addition of V such that it is only present, if present, at regular impurity levels of less than 0.05%, preferably less than 0.02%.
In a further embodiment, the alloys according to this invention have a chemical composition within the ranges of AA7010, AA7040, AA7140, AA7050, AA7081, or AA7085, plus modifications thereof, except they have the higher Si of the present invention in the above-described range of >0.12 to 0.35%, or the higher Si of the present invention in an above-described preferred narrower Si range.
In a preferred embodiment a wrought AA7000-series alloy product according to this invention, consists essentially of, in wt. %:
balance being Al, incidental elements and impurities. Typically such impurities are present each <0.05%, total <0.15%.
In another preferred embodiment a wrought AA7000-series alloy product that can be processed favourable according to this invention, consists essentially of, in wt. %:
The AA7000-series alloy product manufactured according to this invention can be used as an aerospace structural component, amongst others as fuselage sheet, fuselage frame member, upper wing plate, lower wing plate, thick plate for machined parts, thin sheet for stringers, spar member, rib member, floor beam member, and bulkhead member.
In the following, the invention will be explained by the following non-limitative examples.
Two aluminum alloys have been cast having a composition as given in Table 1, and wherein the alloy with 0.02% Si is according to the prior art and the one with 0.23% Si is according to this invention. A regular Ti—C grain refiner was used. The ingots were machined into rolling blocks of 80×80×100 mm. Alloy 1 was given a single homogenisation treatment according to the prior art and that consisted of a controlled heat-up of 30° C./hr from ambient temperature to 470° C. with a 14 hour soak at 470° C. Whereas alloy 2 was given a two-step homogenisation treatment according to the invention that consisted of a controlled heat-up of 30° C./hr from ambient temperature to 470° C. with a 14 hour soak at 470° C., this was followed by a controlled heat-up to 525° C. at 30° C./hr and 7 hours soak. Once the samples had air-cooled, they were preheated to 430° C. and hot-rolled to final gauge of 30 mm. Samples were then solution heat treated at 475° C. with a one-hour soak and then cold water quenched. The samples were then aged to a T76 condition, and subsequently tested for their mechanical properties in three orientations (L, LT, and ST) according to ASTM-E8 standard. The results of which are listed in Table 2, and wherein “TYS” stands for Tensile Yield Strength, “UTS” for Ultimate Tensile Strength and “El” for elongation at fracture. All testing has been done at ½ T.
From the results of Table 2 it can be seen that alloy 2 although having a higher Si content has strength levels better than alloy 1 processed according to prior art practice.
On a pilot scale of testing a billet have been DC-cast having a diameter of 250 mm and a length of over 850 mm. The alloy composition is listed in Table 3, and whereby it is noticed that alloy 3 has an Fe content slightly higher than what is currently customary for aerospace grade rolled products. Alloy 3 would be a typical example of the AA7085 series alloy. From the billet two rolling blocks have been machined having dimensions of 150×150×300 mm. By following this route blocks with an identical chemistry were obtained making it easier to fairly assess the influence of the heat treatments at a later stage on the properties. The blocks were all homogenised using the same cycles of 19 hours at 470° C. whereby industrial heat up rates and cooling rates were applied. Depending on the block a further homogenisation treatment according to the invention was applied whereby the furnace temperature is further increased and where after a second heat treatment or homogenisation treatment of 10 hours at 525° C. was applied. Following the homogenisation the blocks were cooled to room temperature. Thereafter all the blocks were preheated for 5 hours at 450° C. in one batch and hot rolled from 150 to 60 mm. The entrance temperatures (surface measurements) were in the range of 430 to 440° C. and mill exit temperatures varied in the range of 380 to 390° C. After hot rolling the plates received a one or two step solution heat treatment followed by a cold water quench. After a delay of 72 hours the plates were aged to the same T76 temper using a 3-step ageing practice, viz. 6 hours at 120° C., then 12 hours at 154° C. and followed by 24 hours at 120° C. The plates were not stretched prior to ageing. All heat treatments are summarised in Table 4.
The average mechanical properties according to ASTM-B557 standard over 2 samples of the 60 mm plates produced with the various heat treatments are listed in Table 5 and wherein “TYS” stands for Tensile Yield Strength in MPa, UTS for Untimate Tensile Strength in MPa, “El” stands for elongation at fracture in %, and “Kq” for the qualitative fracture toughness in MPa√m. The fracture toughness has been measured in accordance with ASTM B645. The L, LT, L-T and T-L testing was done at ¼ T while ST tensile testing and S-L fracture toughness was done at ½ T.
From the results of Table 5 with respect to the mechanical properties the following can be seen:
Compared to standard processing (Sample 3A1) the variants with a two step treatment according to the invention (Samples 3A2 and 3B2) show a significant increase in toughness, especially in the S-L orientation. It seems that a combined two step homogenisation treatment (Sample 3B2) plus a two step SHT according to this invention provides the best toughness results.
An increase in TYS and UTS is observed for the plates that received a two step SHT (Samples 3A2 and 3B2). However, a two step homogenisation combined a with single step SHT (sample 3B1) gives no improvement. It is not fully clear at the moment, but an assumption is that quenching after SHT from a higher temperature has a positive effect on the ageing responds of Cu containing AA7000-series alloys. Nevertheless, the obtained 20-30 MPa strength increase is considered as an important advantage of the two step SHT according to his invention.
Also the elongation, in particular in ST direction, is significantly improved using the process according to this invention.
Further improvement in toughness can be made by lowering the Fe content to standard aerospace alloy levels.
Sample 3B2 has been tested also for its corrosion resistance in an EXCO test according to ASTM G34, and had a good performance of “EA”.
In a similar approach as with Example 2, two Cu-free 7xxx-series alloys have been produced, the chemical compositions are listed in Table 6. The alloy compositions fall within the compositional range of AA7021. These alloys were processed in a similar approach as with Example 2 and the thermal history is listed in Table 7. The ageing treatment consisted of 24 hours at 120° C. and quenching. The plates were not stretched prior to ageing. The average mechanical properties measured are listed in Table 8.
From the results of Table 8 with respect to the mechanical properties the following can be seen:
Compared to standard processing (Sample 5A1) the variants with a two step treatment according to the invention (Samples 5A2, 5B1, and 5B2) show a significant increase in toughness, especially in the S-L orientation. It seems that a combined two step homogenisation treatment (Sample 5B2) plus a two step SHT according to this invention provides the best toughness results.
The strength is for all variants (5A1 to 5B2) about the same. An increase in ultimate strength and yield strength is not observed in contrast to the results of Example 2 for the Cu containing AA7xxx-series alloys. This result cannot be readily explained.
Compared to the high Si variant (Sample 5A1) the low Si variant (Sample 4A1) the initial toughness values are obviously higher for the low Si alloy composition. However, after two step heat treatment according to this invention the values of the high Si alloy come close to the low Si alloy. The toughness values of the 5B2 sample are still somewhat lower but this is probably due to the fact that 525° C. for the second SHT might just be to low to dissolve all Mg2Si. Employing a higher two step temperature according to the invention would further improve the toughness of the Alloy 5 variants.
Also the elongation, in particular in ST direction, is significantly improved using the process according to this invention.
It is believed that the toughness can be further improved by lowering the Fe content in the aluminum alloy.
Having now fully described the invention, it will be apparent to one of ordinary skill in the art that many changes and modifications can be made without departing from the spirit or scope of the invention as herein described.
This claims the benefit of U.S. provisional application No. 60/818,965, filed Jul. 7, 2006, incorporated herein by reference.
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