WO2009132436A1 - Procédé thermomécanique de traitement d'alliages - Google Patents

Procédé thermomécanique de traitement d'alliages Download PDF

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WO2009132436A1
WO2009132436A1 PCT/CA2009/000560 CA2009000560W WO2009132436A1 WO 2009132436 A1 WO2009132436 A1 WO 2009132436A1 CA 2009000560 W CA2009000560 W CA 2009000560W WO 2009132436 A1 WO2009132436 A1 WO 2009132436A1
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alloy
temperature
recrystallization
heating
sheet
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Shahrzad Esmaeili
David J. Lloyd
Haiou Dr. Jin
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University of Waterloo
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University of Waterloo
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/04—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon

Definitions

  • the present invention relates generally to alloys and processing thereof.
  • the invention relates more specifically to a process for thermomechanically treating alloys to achieve a fine grain structure.
  • Aluminum body panels provide significant weight reductions in automobiles, resulting in fuel efficiency and reduction in greenhouse gas emissions.
  • Al- Mg-Si(-Cu), i.e. AA ⁇ xxx, alloys are used for skin (outer) panel applications for their light weight and high strength achieved through precipitation hardening during paint baking processes, as well as good surface appearance.
  • these alloys have mainly been used for relatively simple shaped panels (e.g. hoods and trunk lids) and/or in high-end car models, they have not been effectively adopted for the production of more complicated shaped panels (e.g. doors) and/or in the mid-range and economy car classes because of their formability issues and high cost.
  • non-heat treatable Al-Mg i.e. 5000 series alloys
  • they are not the alloys of choice for the outer panel applications due to the Ludering problem associated with the magnesium content of the alloys, as well as the lack of hardening potential during paint baking process. Therefore, AA ⁇ xxx alloys remain as the most suitable lightweight choice for outer panels.
  • 6000 series Al sheets in the automotive industry is continuously growing, the limited formability of the current tempers for shaping complex panels, as well as the inherent springback phenomenon associated with room temperature deformation of aluminium alloys has prevented a stronger growth.
  • U.S. Patent No. 6,350,329 (P. Troeger, E.A. Starke, Jr., R. Crooks) [9] describes a fine grain AA ⁇ xxx sheet having a composition close to AA6013 and AA6111, with an average grain size of 10 ⁇ m and superplastic behaviour, processed using conventional cold rolling and a sequence of heat treatments. It is also indicated that the process is suitable for processing any age-hardenable alloy. The processing, characterization, and properties of the fine grain AA ⁇ xxx sheet has also been published as journal papers [10,11].
  • the heat treatment and cold working steps described in U.S. Patent No. ⁇ ,350,329, are as follows: 1. Solution heat treatment followed by rapid cooling to a low temperature
  • the alloy should be cooled after aging preferably by air cooling.
  • Plastic deformation to induce particle simulated nucleation (PSN) of new grains e.g. 20% to induce a total of 80% thickness reduction if Step 2 was ⁇ O%).
  • Static recrystallization preferably using a rapid heat up to the recrystallization temperature to minimize the recovery process (e.g. rapid heating to 54O 0 C , followed by 5 minutes at 54O 0 C).
  • the present invention provides a process of imparting a fine grain structure to a heat treatable alloy, the process comprising: providing a heat treatable alloy having a precipitating constituent; solution heat treating the alloy; cooling the alloy to form a supersaturated solid solution; plastically deforming the alloy to form a dispersed high-energy defect structure for subsequent (a) dispersed nucleation of recrystallization, or (b) dispersed recovery-recrystallization, or both (a) and (b), and to form dispersed fine precipitates; heating the alloy below a recrystallization temperature with a time-temperature profile to continue to form dispersed fine precipitates; and heating the alloy at or above the recrystallization temperature to effect (a) dispersed nucleation of recrystallization and growth of nuclei, or (b) dispersed recovery- recrystallization, or both (a) and (b), to thereby form a fine grain structure.
  • the alloy may be non- isothermally heated to a temperature that is at or above the recrystallization temperature and then may be isothermally heated at the temperature at or above the recrystallization temperature.
  • the heating below the recrystallization temperature and the heating at or above the recrystallization temperature both may consist essentially of non-isothermal heating.
  • the heating below the recrystallization temperature may comprise a mixture of non-isothermal heating and isothermal heating.
  • the heating below the recrystallization temperature may be heating below 300 0 C.
  • the heating at or above the recrystallization temperature may be heating at or above 300 0 C.
  • the fine grain structure may be achieved in the absence of further plastic deformation after the heating below the recrystallization temperature.
  • the alloy may be an aluminum alloy.
  • the alloy may be an AA ⁇ xxx alloy.
  • the alloy may be an AA ⁇ xxx alloy with or without additional incidental or minor alloying elements.
  • the alloy may be a multi- or bi-layered material, wherein at least one layer comprises an AA ⁇ xxx alloy and at least one other layer comprises another aluminum alloy.
  • the another aluminum alloy may be an AAIxxx or AA3xxx alloy.
  • the alloy may be an AA6451 , AA6111 , AA6013, AA6061 , AA6063, or AA6066 alloy.
  • the alloy may be an AA6451 or a Cu-free version of AA6451.
  • the alloy may be a heat treatable titanium, magnesium, cobalt, copper, or nickel alloy.
  • the cooling prior to plastically deforming may comprise quenching.
  • the plastically deforming may comprise cold rolling.
  • the plastically deforming may comprise cold rolling the alloy with a reduction of at least 50%.
  • the process may further comprise naturally aging the alloy.
  • the alloy may be heated continuously to or above the recrystallization temperature. After heating at or above the recrystallization temperature, cooling the alloy or allowing the alloy to cool may be allowed.
  • the alloy is an AA6451 or a Cu-free version of AA6451 ;
  • the plastically deforming comprises cold rolling the alloy with a reduction of at least 50%;
  • the alloy may be non-isothermally heated at a rate of less than 40°C/s to a temperature that is at or above the recrystallization temperature and then is isothermally or non-isothermally heated at the temperature at or above the recrystallization temperature; the average grain size of the fine grain structure is less than 25 ⁇ m; and after heating at or above the recrystallization temperature, the alloy is cooled or allowed to cool.
  • the present invention provides an alloy sheet produced by a process disclosed herein.
  • the alloy sheet may comprise the following properties: a ductility of greater than 80% elongation at 35O 0 C to 55O 0 C and a strain rate of 5x10 4 S "1 to 6.7x10 1 S "1 ; an average grain size of less than 25 ⁇ m; and a grain stability, quantified by an increase in average grain size of less than 100% over one hour at 55O 0 C, and an increase in average grain size of less than 10% over one hour at 400 0 C.
  • the present invention provides an AA ⁇ xxx alloy sheet comprising the following properties: a ductility of greater than 80% elongation at 35O 0 C to 55O 0 C and a strain rate of 5XiO -4 S '1 to 6.7x10 " V 1 ; an average grain size of less than 25 ⁇ m; and a grain stability, quantified by an increase in average grain size of less than 100% over one hour at 55O 0 C, and an increase in average grain size of less than 10% over one hour at 400 0 C.
  • the present invention provides an AA ⁇ xxx alloy sheet comprising the following properties: a ductility of greater than 80% elongation at 35O 0 C to 55O 0 C and a strain rate of 5XiO -4 S "1 to 6.7x10 ' V 1 ; an average grain size of less than 25 ⁇ m; and a substantially uniform precipitate distribution below solvus temperature with precipitates substantially in one of more of the following shapes: oval, rounded, and cuboid.
  • Fig. 1 is a schematic of the process described in U.S. Patent No. 6,350,329;
  • Fig. 2 is a process according to an embodiment disclosed herein (called Route A);
  • Fig. 3 is a process according to an embodiment disclosed herein (called Route B);
  • Fig. 4 is a process according to an embodiment disclosed herein (called Route C);
  • Fig. 5 is a microstructure of an AA6451 sheet processed according to an embodiment disclosed herein (by Route A1 );
  • Fig. 6 is a graph showing the grain size distribution of an AA6451 sheet processed according to an embodiment disclosed herein (Route A1 );
  • Fig. 7 is a microstructure of an AA6451 sheet processed according to an embodiment (Route A1 ), followed by an extended thermal stabilization;
  • Fig. 8 is a graph showing the grain size distribution of an AA6451 sheet processed according to an embodiment disclosed herein (Route A1 ), followed by an extended thermal stabilization;
  • Fig. 9 is an SEM image (back-scattered mode) from an AA6451 sample continuously annealed (at a heating rate of 0.4°C/min) from 5O 0 C to 36O 0 C and held at 36O 0 C for 6 hours;
  • Fig. 10 is an SEM image (secondary electron mode) showing the precipitate structure of a T4P sheet (commercially produced) AA6451 after aging for 2 hours at 35O 0 C;
  • Fig. 11 is an SEM image (secondary electron mode) showing the precipitate structure of a T4P sheet (commercially produced) AA6451 after aging for 2 hours at 35O 0 C;
  • Fig. 12 is a EDS analyzed image of a commercially produced T4P sheet (AA6451 ) after aging for 2 hours at 35O 0 C;
  • Fig 13 is an image of the fine grain (FG) material heated for 2 hours at
  • Fig. 14 is an image of a commercial material (T4P) heated for 2 hours at
  • Fig. 15 is an SEM (secondary electron mode) image showing the precipitate structure of an FG sheet produced by Route A1 , after aging for 2 hours at
  • Fig. 16 is an SEM (secondary electron mode) images showing the precipitate structure of an FG sheet produced by Route A1 , after aging for 2 hours at 350 0 C;
  • Fig. 17 is a high magnification view of a commercial material (T4P) after isothermal heating at 45O 0 C for 2 hours;
  • Fig. 18 is a high magnification view of a FG material after isothermal heating at 45O 0 C for 2 hours;
  • Fig. 19 shows the microstructure of an FG sheet processed through Route
  • A1 (AA6451 );
  • Fig. 20 shows the microstructure of a commercially produced AA6451 sheet
  • Fig. 21 is a graph showing elongation to failure of a commercial material (T4P);
  • Fig. 22 is a graph showing elongation to failure of an FG material
  • Fig. 23 is an SEM fracture surface of a commercial (C) sheet
  • Fig. 24 is an SEM fracture surface of an FG sheet
  • Fig. 25 shows grain size evolution of an FG sheet (AA6451 -ROUTE A1 ) after isothermal hearing at various temperatures, measured using EBSD on planar surface cross-section;
  • Fig. 26 is a graph showing deformation behaviour of an FG sheet and a C sheet at a strain rate of 5.OxIO "4 s "1 ;
  • Fig. 27 is a graph showing deformation behaviour of an FG sheet and a C sheet at a strain rate of 6.7x10 '1 s "1 ;
  • Fig. 28 is a stress strain curve of a T4P sheet tensile tested at 5.0x10 4 s "1 ;
  • Fig. 29 is a stress strain curve of an FG sheet tensile tested at 5.0x10 4 s '1 ;
  • Fig. 30 is a stress strain curve of a T4P sheet tensile tested at 6.7x10 1 s "1 ;
  • Fig. 31 is a stress strain curve of an FG sheet tensile tested at 6.7x10 1 s "1 ;
  • Fig 32 is an optical microscopy image (through-thickness cross section) of deformed (i.e. tensile tested) and fractured sample of a C sheet at 350 0 C and a strain rate of 5.0x10 " V;
  • Fig 33 is an optical microscopy image (through-thickness cross section) of deformed (i.e. tensile tested) and fractured sample of a FG sheet at 350 0 C and a strain rate of 5.0x10 V;
  • Fig. 34 is cross-section of a fractured tensile test sample at 500 0 C and
  • Fig. 35 is a cross-section of a fractured tensile test sample at 500 0 C and
  • thermomechanical process for treating an alloy, for instance an AA6xxx aluminum alloy, to achieve extended high temperature ductility.
  • Such treated alloys may be particularly useful to make automobile panels.
  • the alloy is solution heat treated.
  • the alloy is cooled to form a supersaturated solid solution, and is optionally naturally aged.
  • the alloy is plastically deformed to form a high-energy defect structure for subsequent (a) dispersed nucleation of recrystallization, or (b) dispersed recovery-recrystallization, or both (a) and (b), and to form fine precipitates.
  • the alloy is heated below a recrystallization temperature with a time-temperature profile to continue to form dispersed fine precipitates.
  • thermomechanical processing TMP
  • Figs. 2 to 4 illustrate three of many possible routes of this process
  • Optional storage at about room temperature i.e. during which time natural aging can occur, which can be beneficial
  • a time such as one day, one week, two weeks, or four weeks, which allows for the process to be integrated with existing infrastructure and operations
  • sufficient plastic deformation e.g. at least 30%, or at least 50%, or 60 to 80% reduction, dependent on the alloy type, by conventional or asymmetric cold rolling.
  • This induces a high energy defect structure (large dislocation density), with a fine distribution of small solute clusters depending on the storage time at about room temperature.
  • the minimum thickness reduction will depend on the alloy. There is no specific maximum thickness reduction.
  • the preferred reduction amount based on current commercial line capabilities for the initial thickness and common aluminum sheet thickness for automotive panels, for AA ⁇ xxx alloys, is about 60 to 70%; for instance beginning with a 2.5mm sheet and finishing with a 1mm sheet, or beginning with a 5mm plate and finishing a 1.5mm sheet. These thicknesses are mere examples. Various starting and finishing thicknesses may be selected, and may be based on, for instance, the alloy, the desired product specification, or line requirements. [0047] 3. Annealing through heating of the as-deformed material from a low temperature (e.g., room temperature, 50°, 100 0 C, less than 300 0 C, or less than the recrystallization temperature of the alloy) to a high enough temperature (e.g.
  • a low temperature e.g., room temperature, 50°, 100 0 C, less than 300 0 C, or less than the recrystallization temperature of the alloy
  • a high enough temperature e.g.
  • the lowest preferred starting temperature may depend on the capability of the infrastructure/production line or the preferred operating conditions, as well as the alloy type.
  • the maximum starting temperature may be between 250 to 35O 0 C for AA6xxx alloys, corresponding to the approximate temperature range for the "nose" (also known as
  • the starting temperature lower than the nose of the TTT diagram will increase the number density of precipitates within the grains and decrease the possibility of grain boundary precipitation due to a high driving force for homogeneous nucleation of precipitates and/or precipitation on dislocations accumulated during deformation.
  • a high number density of precipitates within the grains and prevention of any major grain boundary precipitation are factors in creating the desirable sheet structure.
  • precipitates may include all precipitate forms such as solute clusters, GP zones, metastable precipitates, transitional precipitates, and equilibrium precipitates.
  • the finishing temperature is a temperature at which a substantial fraction of grains recrystallize (e.g.
  • fraction recrystallized 90%, 95% or 99%
  • a shorter required time usually requires a higher temperature, as recrystallization is a thermally-activated process and therefore dependent on time at temperature.
  • a higher % reduction in the thickness during the deformation stage can also significantly affect the maximum temperature achieved.
  • a higher percent reduction in the thickness usually requires lower temperature and/or shorter times to achieve substantial recrystallization.
  • the preferred maximum temperature is ⁇ 30O 0 C in AA ⁇ xxx alloys.
  • the maximum limiting temperature is the practical solidus temperature (i.e. melting should not happen).
  • the preferred annealing process is ramp heating from the selected low temperature with a selected heating rate (e.g. 0.4°C/min). This is preferred to make use of the existing aluminum rolled product production lines with minimal requirement for change in the production line and no wait time or batch annealing process. There is no specific minimum heating rate and the maximum will depend on the particular alloy used and whether the ramp heating is combined with any other heating and/or cooling step (e.g. isothermal step or another non-isothermal heating, or a cooling step).
  • a selected heating rate e.g. 0.4°C/min
  • the maximum ramp heating rate is also dependent on the choice of the maximum temperature; a lower maximum rate is needed when the highest temperature is relatively high (e.g. 0.4 0 C /min to reach 38O 0 C and no isothermal time at this temperature, or 100°C/min to reach 31O 0 C and 24 hours at this temperature). If, after the initial ramp heating, there is at least one additional step which provides some time below the nose of the TTT diagram, the ramp heating can be as fast as possible. (As an example for AA ⁇ xxx alloys, extremely fast heating from room temperature to 100 0 C, 2 hours at 100 0 C 1 4 0 C /min to 36O 0 C, 1 hour at 36O 0 C).
  • the key requirement is to allow sufficient time below the nose of the TTT diagram to form a high number density of fine precipitates and then allow enough time at a relatively high temperature to achieve the recrystallization requirement.
  • the large number density of fine precipitates either prevents recovery and the rapid formation of dislocation cell structures or it promotes a continuous recovery-recrystallization process.
  • the inhibition of recovery preserves a large driving force for the nucleation of new deformation-free grains (i.e. nucleation of recrystallization) at all, or most, high energy sites (e.g. abundant dislocation bundles, shear bands, around large pre-existing intermetallic particles) and therefore results in a fine grain microstructure.
  • U.S. Patent No. 6,350,329 uses, after solution the heat treating, the plastic deformation, and the annealing, another plastic deformation and subsequent static recrystallization.
  • a second deformation is required.
  • a second deformation is not required as the first deformation is sufficient to produce the defect structure required by the subsequent step.
  • fine precipitates form (such as nano-size). They continue to form during the early stage of annealing (called nucleation and growth of precipitates).
  • the nuclei can be in the form of atom clusters or well developed zones, precipitates, etc. Some time at relatively low temperatures allow the disperse formation of these precipitates.
  • the precipitate nucleation and growth is usually enhanced due to the defect structure (high energy state promotes nucleation and the defects promote diffusion (i.e. movement) of solute atoms which attach themselves to the nucleated precipitates and cause their growth to be faster). These finely distributed precipitates then pin the dislocation motion so that their recovery (i.e. elimination) does not happen quickly and they remain as high energy sites.
  • Routes A and B show preferred annealing routes for their enhanced compatibility with current commercial sheet production lines. It is noted that only one stage of deformation is present (e.g. rolling).
  • Route C shows annealing where the heating is not solely non- isothermal. That is, there may be one or more steps of isothermal heating.
  • Alternative routes may be used in accordance with the production line requirements or preferred operations or final product properties (e.g. fraction recrystallized, level of strength, ductility or potential for better corrosion resistance). These routes can have multiple isothermal and non-isothermal steps including heating and cooling cycles, heating with non-linear rates, and non-linear rate heating and cooling cycles.
  • the process may provide a cost-effective and efficient process of sheet production with the use of existing production facilities and minimal need for change in the production lines.
  • the produced sheets of embodiments of the instant invention may be a better choice for the automotive industry to achieve cost- effective forming of complex shape panels, which require high levels of ductility.
  • the resultant alloy sheets could be used in various applications where one or more properties of these resultant alloy sheets, for instance the fine grain size and good ductility, could be exploited.
  • Heat treatable alloys such as a 2xxx, 6xxx, 7xxx and some 8xxx aluminum alloys having a precipitating constituent, are conceivably candidates for processing in accordance with embodiments of this invention.
  • an alloy In order for an alloy to be heat treatable, it should have at least one alloying element that is different from the element of the matrix, (e.g. Al-Cu alloys with Al acting as the matrix and Cu acting as the precipitating constituent forming a series of Cu-Al containing precipitate phases).
  • the precipitates are used during annealing to help create and retain a fine grain structure by affecting recovery, recrystallization, and grain growth processes.
  • the alloy is cooled to form a supersaturated solid solution.
  • the mode of cooling is not critical, rapidly cooling the alloy to a temperature at which the diffusion rate of any of the elements in the alloy is not appreciable, and the formation of precipitates, particularly on the grain boundaries, is practically prevented, ensures the retention of as much solute in solid solution as possible, making the maximum amount of solute available for the subsequent formation of finely distributed precipitates during early stage annealing.
  • Some precipitation, such as solute cluster formation may occur during rapid cooling, and is acceptable. Prevention of major precipitate formation also ensures low strength and therefore ease of deformation during the plastic deformation step.
  • the rapid cooling may be accomplished, for example, by quenching in a medium such as water, oil or air, or another known rapid cooling mechanism.
  • the alloy is sufficiently plastically deformed to produce a high-energy defect structure useful for subsequent (a) dispersed nucleation of recrystallization and growth of nuclei, or (b) dispersed recovery- recrystallization, or both (a) and (b).
  • the deformation of the solution heat-treated alloy is preferably carried out at room temperature, although this temperature will vary with alloy composition. This step also may be carried out at other temperatures. Most preferably, the deformation is performed at whatever temperature is most convenient and economical, provided that sufficient energy is retained in the alloy for the formation of a high-energy defect structure.
  • a high-energy defect structure is that magnesium alloys might be deformed at a higher temperature to facilitate extensive reduction in thickness, not achievable at room temperature due to the hep structure of magnesium.
  • some alloying elements enhance the work hardening behavior of alloys when such alloying elements are present in solid solution. For example, magnesium is known to have this effect in aluminum alloys.
  • aluminum alloys containing Mg in solid solution may develop greater stored strain energy for a given amount of deformation than alloys not containing Mg (some alloying additions, such as Mg in solid solution in Al matrix are known to lower the dynamic recovery rate). Accordingly, the high-energy defect structure required for such alloys may be more readily attainable for alloys containing one or more alloying elements in solution than for alloys not having such alloying elements.
  • the potential solute cluster formation during rapid cooling or optional storage can also retain or enhance the work hardening behavior of the solid solution alloy, and thus the level of energy storage during deformation, in some alloys such as AA ⁇ xxx alloys. These clusters can act as further heterogeneous nucleation sites (or practically as nuclei) for precipitate formation during annealing.
  • the alloy is heated to form fine substantially homogenous distribution of precipitates during early stage (i.e. low temperature) annealing. Again, this is referred to herein as dispersed fine precipitates and may include some areas of heterogeneity.
  • the preferred times and temperatures for the heating process are dependent upon the type of alloy used.
  • These finely distributed precipitates grow and/or coarsen during further annealing (at higher temperature) such that they interact with recovery and recrystallization mechanisms and induce the formation of equi-axed, or relatively equi-axed, new grains with a fine average grain size and a relatively narrow size distribution.
  • the precipitates can (or tend to) later inhibit grain growth during extended exposure to high temperatures.
  • the cooling rate after annealing is not critical if the final annealing temperature is below the solvus temperature. If it is above the solvus temperature, a relatively fast cooling rate (e.g. air cooling or cooling with/in a liquid medium at a relatively low temperature such as water at room temperature) is preferred to retain the solutes in the solid solution matrix and mitigate grain boundary precipitation.
  • a relatively fast cooling rate e.g. air cooling or cooling with/in a liquid medium at a relatively low temperature such as water at room temperature
  • Microstructure A fine grain sheet with a relatively equi-axed grain structure was obtained by a process according to a disclosed embodiment, using continuous annealing including non-isothermal heating from 5O 0 C to 38O 0 C for 13 hours and isothermal stabilization at 380 0 C for 20 minutes (called hereafter Route A1 as it falls within Route A of Fig. 1 ), applied on AA6451.
  • Route A1 as it falls within Route A of Fig. 1
  • the microstructure of this sheet and the grain size distribution are shown in Figs. 5 and 6, respectively.
  • the average grain size of the sheet (through thickness) is approximately 8.0 ⁇ m (approximately 10 to 11 microns in sheet plane).
  • the microstructure has excellent stability at high temperature. This is evidenced by examining the microstructure of the alloy when the stabilizing step at 38O 0 C is extended to 5 hours. The result is shown in Figs. 7 and 8. The microstructure has barely changed when the alloy is exposed to a temperature of 380 0 C for an extended period of time.
  • Fig. 25 shows grain size evolution of the FG sheet (AA6451 -ROUTE A1 ) after isothermal heating at various temperatures, measured using EBSD on planar surface cross-section. Grain size stability is observed.
  • the significance of the relatively stable microstructure is that the grains will resist growing during high temperature deformation and therefore will remain fine for relatively long times at high temperatures. This will result in extended ductility (formability) at high temperatures. If the grains would tend to grow quickly, the microstructure would not remain fine-grained during high temperature deformation. It would then render the material coarse-grained one, similar in grain size and/or size distribution to the existing commercial sheets, with potentially inferior ductility at high temperatures and potentially prone to premature failure during high temperature deformation.
  • the precipitate morphology (more rounded) and distribution (substantially uniformly distributed everywhere) will make the sheet more ductile compared to the current commercial coarse-grained sheet. Further, the sheet will also still be less prone to premature failure compared to the commercial sheet because of the precipitate structure difference.
  • Another characteristic of the microstructure is the precipitate distribution.
  • Fig. 9 shows an SEM image (back-scattered mode) from an AA6451 sample continuously annealed (at a heating rate of 0.4°C/min) from 50 0 C to 36O 0 C for 6 hours. Precipitates are observed as black particles when back scattered electrons are used to obtain the SEM image.
  • the precipitate distribution is very different in the microstructure of the commercially-produced sheets when exposed to high temperatures, particularly for an extended period of time.
  • Figs. 10 and 11 are SEM images (secondary electron mode) showing the precipitate structure of a T4P sheet (commercially produced) AA6451 after aging for 2 hours at 35O 0 C. A network of grain boundary precipitates and fine distribution of small, elongated particles are seen within the grains.
  • Fig. 12 is an EDS analyzed image of a commercially produced T4P sheet
  • Figs. 13 (FG sheet) and 14 (T4P) at 35O 0 C show that for the T4P material, grain boundary precipitates can be both small and round and elongated.
  • Figs. 15 and 16 are SEM (secondary electron mode) images showing the precipitate structure of an FG sheet produced by Route A1 , after aging for 2 hours at 35O 0 C.
  • Fig. 17 is a high magnification view of the T4P material after isothermal heating at 45O 0 C for 2 hours.
  • Fig. 18 is a high magnification view of the FG material after isothermal heating at 450 0 C for 2 hours.
  • Figs. 10 to 18 demonstrate the effect of high temperature exposure to precipitate distribution in the microstructure of a commercially-produced AA6451 sheet.
  • Precipitates are white and light gray particles within the grains and along the grain boundaries (white and light grey network).
  • Figures 10 and 11 after aging for 2 hours at 350 0 C, large, blocky and elongated precipitates grow on the grain boundaries, which were not seen in the as-received material.
  • the interior of the grains can be seen to contain smaller, more homogeneously distributed thin elongated precipitates (see Fig. 12).
  • the precipitates at grain boundaries and in grain interiors grow preferentially in certain directions (see Fig.
  • the T4P sheet shows precipitates both inside the grains (finely distributed) and along grain boundaries, which form a pronounced network of precipitates; and areas adjacent to the grain boundaries do not have the finely distributed precipitates, as within the interior of the grains. These areas, which are defined as precipitate free zones (PFZ), are weak and prone to premature failure during deformation.
  • PFZ precipitate free zones
  • the precipitate network itself is also considered deleterious to the properties and/or the performance of the alloys, because of effects such as inducing brittle fracture along the grain boundaries, or stress concentration effects due to the closely spaced elongated precipitates.
  • the precipitate network also creates corrosion problems.
  • Figs. 19 and 20 provide another direct comparison of the microstructure of the fine-grained sheet produced through Route A1 when exposed to a high temperature with that of the commercial sheet exposed to the same heating condition. The difference in the distribution of precipitates in the two sheets is noticeable.
  • the sheet of Figs. 19 and 20 provide another direct comparison of the microstructure of the fine-grained sheet produced through Route A1 when exposed to a high temperature with that of the commercial sheet exposed to the same heating condition. The difference in the distribution of precipitates in the two sheets is noticeable.
  • Figure 19 and 20 are AA6xxx alloy sheets (thermomechanically processed and then heated for 15 minutes at 45O 0 C.
  • Figure 19 is a fine grained sheet processed through Route A1 (AA6451) and
  • Figure 20 is a commercially produced AA6451 sheet.
  • Fig. 22 summarizes the total percent elongation to failure of the FG material at each temperature and strain rate.
  • Figs. 26 and 27 show deformation behaviour of an FG sheet and a C sheet deformed at 350°C and strain rates of 5.0x10 4 s "1 (Fig. 26) and 6.7x10 1 s '1 .
  • Figs. 28 and 29 show deformation behaviour of FG sheets and C sheets at various temperatures. It is evident from Figures 18 through 20 that except in three cases of test temperatures and strain rates, the FG sheet achieves significantly higher % elongations at high temperatures, with more pronounced difference at 35O 0 C and 40O 0 C (i.e. temperatures of interest for warm forming operations).
  • the two materials showed similar % elongation at 500°C and the strain rates of 6.7x10 "2 s “1 and 6.7x10 "1 s “1 and at 550°C and the strain rate of 6.7x10 "1 s “1 . These results reflect the FG sheet's high potential for future cost-effective warm forming operations in a wide range of strain rates and temperatures.
  • the highest % elongation of approximately 290% (measured using a video extensometer) achieved at 500 0 C and strain rate of 2.0x10 2 s "1 (this is a much higher strain rate than the one reported for the superplastic sheet formed in U.S. Patent No. 6,350,329, mentioned above (i.e.
  • Figs. 32 and 33 show optical microscopy images (through-thickness cross section) of deformed (i.e. tensile tested) and fractured samples of the FG sheet (Fig. 33) and a C sheet (Fig. 32) at 350 0 C and a strain rate of 5.0x10 " V 1 .
  • the FG materials have shown more effective and uniform thinning to the fracture surface (which appears as a point when viewing through-thickness cross section).
  • the C sheet has a rugged through- thickness fracture profile, demonstrating the inferior deformation to fracture behaviour.
  • the FG material has also shown, generally, a higher resistance to cavitation during deformation, particularly when deformed at low strain rates and 500 0 C. This is demonstrated in Figs.
  • Another difference in the properties of the FG sheet with the C sheet is in achieving large % elongations without encountering premature failure due to the effect of precipitate network at the grain boundaries and the associated precipitate-free zones (PFZ).
  • PFZ precipitate-free zones
  • embodiments of the instant invention include alloy sheets with one or more of the following properties: average grain size: ⁇ 40 ⁇ m, ⁇ 25 ⁇ m, ⁇ 20 ⁇ m, ⁇ 18 ⁇ m, ⁇ 16 ⁇ m, ⁇ 14 ⁇ m, or ⁇ 12 ⁇ m; and ductility according to one or any combination of the elongation % values in Table ! [0098] Table 1 : Elongation values at particular strain rates and temperatures.
  • the comparable commercial alloy of AA6111 has an average grain size in the range of 30 to 45 microns. Thus, this thermomechanical process has resulted in grain refining.

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Abstract

La présente invention concerne un procédé thermomécanique de traitement d'un alliage, par exemple un alliage d'aluminium AA6xxx, permettant d'obtenir une plus grande ductilité à haute température. De tels alliages traités peuvent être particulièrement utiles pour fabriquer des panneaux d'automobiles. Au cours d'une première étape, l'alliage subit un traitement thermique de mise en solution. Puis l'alliage est refroidi de façon à former une solution solide sursaturée et il est éventuellement soumis à un vieillissement naturel. L'alliage est ensuite déformé plastiquement de manière à former une struture lacunaire à haute énergie permettant par la suite (a) une nucléation dispersée de recristallisation, ou (b) une recristallisation à récupération dispersée, ou les deux (a) et (b), et à former de fins précipités dispersés. L'alliage est alors chauffé à une température inférieure à une température de recristallisation selon un profil temps-température permettant de continuer à former les fins précipités dispersés. Puis l'alliage est chauffé jusqu'à une température égale ou supérieure à la température de recristallisation de façon à procéder à (a) une nucléation dispersée de recristallisation et une croissance des noyaux, ou à (b) une recristallisation à récupération dispersée, ou les deux (a) et (b), ce qui permet d'obtenir une structure à grain fin. On peut ensuite refroidir ou laisser refroidir l'alliage.
PCT/CA2009/000560 2008-04-28 2009-04-28 Procédé thermomécanique de traitement d'alliages Ceased WO2009132436A1 (fr)

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Cited By (10)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
FR2979354A1 (fr) * 2011-08-31 2013-03-01 Peugeot Citroen Automobiles Sa Procede de traitement d'une piece en alliage d'aluminium
US8999079B2 (en) 2010-09-08 2015-04-07 Alcoa, Inc. 6xxx aluminum alloys, and methods for producing the same
US9469892B2 (en) 2010-10-11 2016-10-18 Engineered Performance Materials Company, Llc Hot thermo-mechanical processing of heat-treatable aluminum alloys
US9587298B2 (en) 2013-02-19 2017-03-07 Arconic Inc. Heat treatable aluminum alloys having magnesium and zinc and methods for producing the same
US9926620B2 (en) 2012-03-07 2018-03-27 Arconic Inc. 2xxx aluminum alloys, and methods for producing the same
US10301709B2 (en) 2015-05-08 2019-05-28 Novelis Inc. Shock heat treatment of aluminum alloy articles
US10570490B2 (en) 2015-04-08 2020-02-25 Baoshan Iron & Steel Co., Ltd. Strain-induced age strengthening in dilute magnesium alloy sheets
CN114364470A (zh) * 2019-09-04 2022-04-15 赛峰飞机发动机公司 限制工件中出现再结晶晶粒的制造金属工件的方法
CN115852276A (zh) * 2022-12-23 2023-03-28 中国科学院金属研究所 一种抑制面心立方高熵合金中有害晶界析出相析出的晶界调控工艺
US11874063B2 (en) 2016-10-17 2024-01-16 Novelis Inc. Metal sheet with tailored properties

Citations (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4092181A (en) * 1977-04-25 1978-05-30 Rockwell International Corporation Method of imparting a fine grain structure to aluminum alloys having precipitating constituents
US4722754A (en) * 1986-09-10 1988-02-02 Rockwell International Corporation Superplastically formable aluminum alloy and composite material
US4797164A (en) * 1986-09-30 1989-01-10 Swiss Aluminum Ltd. Process for manufacturing a fine-grained recrystallized sheet
US6350329B1 (en) * 1998-06-15 2002-02-26 Lillianne P. Troeger Method of producing superplastic alloys and superplastic alloys produced by the method

Patent Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4092181A (en) * 1977-04-25 1978-05-30 Rockwell International Corporation Method of imparting a fine grain structure to aluminum alloys having precipitating constituents
US4092181B1 (fr) * 1977-04-25 1985-01-01
US4722754A (en) * 1986-09-10 1988-02-02 Rockwell International Corporation Superplastically formable aluminum alloy and composite material
US4797164A (en) * 1986-09-30 1989-01-10 Swiss Aluminum Ltd. Process for manufacturing a fine-grained recrystallized sheet
US6350329B1 (en) * 1998-06-15 2002-02-26 Lillianne P. Troeger Method of producing superplastic alloys and superplastic alloys produced by the method

Cited By (15)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US8999079B2 (en) 2010-09-08 2015-04-07 Alcoa, Inc. 6xxx aluminum alloys, and methods for producing the same
US9194028B2 (en) 2010-09-08 2015-11-24 Alcoa Inc. 2xxx aluminum alloys, and methods for producing the same
US9249484B2 (en) 2010-09-08 2016-02-02 Alcoa Inc. 7XXX aluminum alloys, and methods for producing the same
US9359660B2 (en) 2010-09-08 2016-06-07 Alcoa Inc. 6XXX aluminum alloys, and methods for producing the same
US9469892B2 (en) 2010-10-11 2016-10-18 Engineered Performance Materials Company, Llc Hot thermo-mechanical processing of heat-treatable aluminum alloys
EP2627794A4 (fr) * 2010-10-11 2017-12-27 Engineered Performance Materials Company LLC Traitement thermo-mécanique à chaud d'alliages d'aluminium de traitement thermique
FR2979354A1 (fr) * 2011-08-31 2013-03-01 Peugeot Citroen Automobiles Sa Procede de traitement d'une piece en alliage d'aluminium
US9926620B2 (en) 2012-03-07 2018-03-27 Arconic Inc. 2xxx aluminum alloys, and methods for producing the same
US9587298B2 (en) 2013-02-19 2017-03-07 Arconic Inc. Heat treatable aluminum alloys having magnesium and zinc and methods for producing the same
US10570490B2 (en) 2015-04-08 2020-02-25 Baoshan Iron & Steel Co., Ltd. Strain-induced age strengthening in dilute magnesium alloy sheets
US10301709B2 (en) 2015-05-08 2019-05-28 Novelis Inc. Shock heat treatment of aluminum alloy articles
US11874063B2 (en) 2016-10-17 2024-01-16 Novelis Inc. Metal sheet with tailored properties
CN114364470A (zh) * 2019-09-04 2022-04-15 赛峰飞机发动机公司 限制工件中出现再结晶晶粒的制造金属工件的方法
CN114364470B (zh) * 2019-09-04 2023-08-04 赛峰飞机发动机公司 限制工件中出现再结晶晶粒的制造金属工件的方法
CN115852276A (zh) * 2022-12-23 2023-03-28 中国科学院金属研究所 一种抑制面心立方高熵合金中有害晶界析出相析出的晶界调控工艺

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