EP3456853B1 - Manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates - Google Patents

Manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates Download PDF

Info

Publication number
EP3456853B1
EP3456853B1 EP17468002.5A EP17468002A EP3456853B1 EP 3456853 B1 EP3456853 B1 EP 3456853B1 EP 17468002 A EP17468002 A EP 17468002A EP 3456853 B1 EP3456853 B1 EP 3456853B1
Authority
EP
European Patent Office
Prior art keywords
precipitates
mass
ageing
alloys
alloy
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Active
Application number
EP17468002.5A
Other languages
German (de)
French (fr)
Other versions
EP3456853A1 (en
Inventor
Franc Zupanic
Tonica Boncina
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Univerza v Mariboru
Original Assignee
Univerza v Mariboru
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Univerza v Mariboru filed Critical Univerza v Mariboru
Publication of EP3456853A1 publication Critical patent/EP3456853A1/en
Application granted granted Critical
Publication of EP3456853B1 publication Critical patent/EP3456853B1/en
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Classifications

    • 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
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/12—Alloys based on aluminium with copper as the next major constituent
    • 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
    • C22F1/002—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working by rapid cooling or quenching; cooling agents used therefor
    • 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
    • C22F1/057—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 of alloys with copper as the next major constituent

Definitions

  • the invention belongs to the area of non-ferrous metallurgy, more specifically to the casting and heat treatment of aluminium alloys.
  • the aluminium and aluminium alloys are characterized by low densities, high specific strengths (strength divided by density), and competitive corrosion resistance. Aluminium parts can be produced economically using versatile manufacturing technologies. By using standard strengthening mechanisms, limiting values of tensile strengths between 550 MPa and 600 MPa can be achieved in wrought alloys, but with reduced elongation (around 10 %). In Al-Si casting alloys alloyed with different alloying additions (e.g. AlSi12CuNiMg), tensile strengths up to 400 MPa can be attained, and the elongation is around 5 %.
  • different alloying additions e.g. AlSi12CuNiMg
  • the most typical heat treatment is designed by T6.
  • the T6 heat treatment consists of solution treatment at temperatures above the alloy's solvus temperature. This treatment results in homogenising the alloy and transferring the precipitates-forming alloying elements from different intermetallic phases into an Al-rich solid solution. Alloy is then quenched. The cooling rate is to be sufficient to retain almost all dissolved alloying elements in the solid solution.
  • natural ageing at room temperature
  • artificial ageing at elevated temperatures, typically below 200 °C
  • the alloying elements start to precipitate from the solid solution. The precipitates strengthen the aluminium alloy.
  • Heat treatment T5 is an alternative to T6.
  • the alloy is directly heated from temperature to the temperature of artificial ageing without solution treatment.
  • the prerequisite for this is, for example by casting, that sufficiently high-cooling rates are achieved, which retain a high content of alloying elements in the solid solution. Precipitates are formed during direct ageing.
  • aluminium alloys should be manufactured in different shapes, having an appropriate combination of low density, high-strength at room temperature, heat resistance, and competitive price.
  • High-strength and heat-resistant precipitation hardening alloys without Sc loose their heat resistance at temperatures slightly above 200 °C because the precipitates become coarser or even dissolve themselves.
  • Precipitation-hardened alloys with Sc have higher heat resistance, but are more expensive because of Sc.
  • the content of Sc in these alloys is rather high. Because of scarcity of Sc, it is not possible to produce alloys in higher amounts.
  • the alloys, strengthened only by quasicrystals, can be produced by rapid solidification; therefore, their use is rather limited.
  • High-strength aluminium alloys are based mainly on heat-treatable alloys from the systems Al-Cu (series AA2xxx), Al-Si-Mg (series AA6xxx), Al-Li and Al-Zn-Mg (series AA7xxx).
  • the alloys are produced by different techniques (casting, powder metallurgy, forming), and then heat-treated mainly by T6 and T5 heat treatments.
  • Increased heat resistance is achieved by the addition of Mn, Cr, Zr, Nb, Hf or V, which can form different types of precipitates that are less soluble in the aluminium matrix, and can remain in the alloy at elevated temperatures.
  • the patent WO 2002063059 A1 discloses a two-stage ageing process of an article made from an alloy comprising at least aluminium and copper.
  • the patent WO 2008003503 A2 discloses the method of manufacturing of AA2xxx alloys for manufacturing aluminium wrought products in relatively thick gauges. In both alloys, different Al-Cu precipitates are formed during manufacturing that cannot withstand temperatures above 200 °C.
  • the patent US 5759302 A describes the compositions of Al-Cu and Al-Zn-Mg alloys, alloyed with Mn, Cr or Zr, and heat treating procedures to obtain dispersoids from the systems Al-Mn, Al-Cr or Al-Zr that can contribute to increased fracture toughness, fatigue resistance and formability.
  • the patent US 20170051383 A1 determines the composition of the alloys and parameters of the three-step ageing process to obtain dual precipitates in the aluminium matrix that can provide the application of the said alloys for high-temperature applications.
  • the alloy is alloyed with Cu and Zr, and with Nb, Hf or V.
  • the patent US 5226983 A discloses the Al-Zr-Li-X alloys that are produced by rapid solidification. The alloy is then subjected to multiple ageing treatments after solution treatment. The microstructure of the alloy is composed of aluminium matrix and Al 3 (Zr, Li) precipitates.
  • the patent WO 2011122958 A1 optimises the Al-Mg-Si-Cu alloy for high-temperature stability. It defines the chemical compositions of the alloy. The aim is to obtain an L-phase as a dominant precipitate type about its number density in the over-aged state.
  • the patent US 6074498 A defines Al-Cu-Li-Sc alloys, which are subjected to a dual ageing treatment. The purpose is to attain an array of fine T1 phase precipitates within the aluminium grain, leaving a substantially T1 phase precipitate-free zone along the grain boundaries, and an array of coarse ⁇ ' and ⁇ ' phase precipitates throughout the grains with little or no ⁇ ' and ⁇ ' phase-free zones.
  • L 12 -precipitates have an ordered cubic crystal structure with a general formula Al 3 X, where X presents one or more elements, the most important are Sc, Zr, and Y.
  • the L 12 -precipitates are coherent with Al-rich solid solution. They are formed during heat treatment, and are rather stable at elevated temperatures.
  • Table 1 gives a list of patents for Al-alloy containing L 12 -precipitates. The patents disclose the composition, manufacturing procedures, e.g. casting, rapid solidification, rolling, welding, extrusion, as well as their heat treatment of aluminium alloys.
  • the patent US 6248453 B1 discloses the aluminium alloys with Sc, Er, Lu, Yb, Tm and U, and with at least one element selected from the group consisting of Mg, Ag, Zn, Li, and Cu.
  • the alloy produced with the said composition and rapid solidification comprises of Al-matrix and Al 3 X precipitates with the L 12 -structure.
  • the patent US 5620652 A discloses the aluminium alloys containing scandium with zirconium additions that can be used for recreational products, ground transportation, marine and aerospace structures. No details of heat treatment are given. It can be supposed that they predominantly consist of Al-rich matrix and Al 3 X precipitates with L 12 -structures.
  • the patent US 20100143185 A1 discloses a method for producing high-strength aluminium alloy powder containing L 12 intermetallic dispersoids.
  • the powder is produced by melt atomization.
  • the patented compositions of alloys cause the formation of L 12 -precipitates in the powder, which are uniformly distributed in Al-rich matrix.
  • Table 1 List of patents for Al-alloys containing L 12 precipitates EP 2598664 B1 HIGH TEMPERATURE SCANDIUM-CONTAINING ALUMINIUM ALLOY WITH IMPROVED EXTRUDABILITY RU 2 446 222 HEAT-RESISTANT ALUMINIUM-BASED ALLOY AND METHOD FOR OBTAINING DEFORMED SEMI-FINISHED PRODUCTS FROM IT RU 2 287 600 ALUMINUM-BASE MATERIAL US7811395 , US7871477 , US7909947 , US8002912 , US20090260722 , US20090260723 , US20090263273 High strength L12 aluminum alloys US7875131 L12 strengthened amorphous aluminum alloys US7875133 , US20090260724 , US20090260725 , US20110041963 , US7883590 Heat treatable L12 aluminum alloys US7879162 High strength aluminum alloys with L12 precipitates US80
  • the patents RU 2 446 222 and RU 2 287 600 deal with deformable aluminium alloys, with scandium, zirconium, manganese, copper and vanadium.
  • the comparison of the patented chemical composition with the chemical composition of the present application is given in Table 2.
  • the composition ranges for copper and manganese do not overlap.
  • the content of zirconium in the current application is smaller than in the Russian patents.
  • zirconium can be replaced by the other elements (Y, Ti, Hf, or Nb).
  • the alloys in the patents RU 2 446 222 and RU 2 287 600 do not contain beryllium, while in this application, the necessary content of Be lies between 0.001 and 0.9 mass.% Be.
  • Al-Cu-Mn-Zr-Sc alloy which does not require solution treatment and quenching for obtaining high-strength.
  • This alloy contains L 12 precipitates, but do not have icosahedral quasicrystalline precipitates. Instead of IQC-precipitates, Al 20 Mn 3 Cu 2 precipitates formed.
  • Their disadvantage is that they are in the form of rods. In addition, their number density is much lower than the density of icosahedral precipitates with spherical morphology.
  • Quasicrystals are crystals having long-range order, but without periodicity. This is the main difference with periodic crystals.
  • Aluminium quasicrystal-strengthened alloys possess periodic aluminium-rich matrix, in which a quasicrystalline phase is dispersed.
  • the quasicrystalline phase can form during rapid solidification and casting, or it is added to aluminium by methods typical for manufacturing composites.
  • the quasicrystalline particles can be added to an aluminium melt; parts are then obtained by casting.
  • the quasicrystalline powder is added to aluminium powder, and products are produced via powder metallurgical routes.
  • the patent US 5593515 A discloses the manufacturing of alloy with a general formula Al bal Q a M b X c T d , wherein Q represents at least one element selected from the group consisting of Mn, Cr, V, Mo and W; M represents at least one element chosen from the group consisting of Co, Ni, Cu and Fe; X represents at least one element selected from rare earth elements including Y or Mm; T represents at least one element selected from the group consisting of Ti, Zr and Hf; and a, b, c and d represent the following atomic percentages: 1 ⁇ a ⁇ 7, 0>5, 0>c ⁇ 5 and 0>d ⁇ 2.
  • the alloy of this invention is said to be excellent in the hardness and strength at both room temperature and a high temperature, and also in thermal resistance and ductility.
  • the products can be of a rapidly solidified alloy, a heat-treated alloy obtained by heat-treating a rapidly solidified alloy, and a compacted and consolidated alloy material obtained by compacting and consolidating a rapidly solidified alloy.
  • the patent US 5419789 A discloses the aluminium-based alloy which consists Al and 0.1 to 25 atomic % of at least two transition metal elements and has a structure in which at least quasicrystals are homogeneously dispersed in a matrix composed of Al or a supersaturated Al solid solution.
  • the alloy is in the form of a rapidly solidified material, a heat treated material of the rapidly solidified material, or a compacted and consolidated material formed from the rapidly solidified material.
  • the milled and extruded melt-spun ribbons of the alloy Al 92 Mn 6 Ce 2 attained strength of approximately 800 MPa and elongation ⁇ 25 %, but the conventionally cast rods had strength around 500 MPa and elongation ⁇ 20 %.
  • Kim et al. [6] found decagonal quasicrystalline precipitates in as-rapidly solidified and annealed commercial AISI 2024 aluminium alloy containing 2 wt% Li.
  • Zhang et al. [7, 8] studied the annealing of supersaturated solid solutions in some rapidly solidified Al-transition metal alloys using transmission electron microscopy. They found that icosahedral quasicrystalline precipitates formed in the Al-matrix in the binary systems Al-Cr and Al-Fe, and also in the systems Al-V, Al-Cr and Al-Mo, when they were alloyed by a small amount of Fe.
  • Zupanic et al. [9] discovered icosahedral precipitates in an Al-Mn-Be-Cu that was cast into a copper mould, after 24 h of annealing at 300 °C.
  • the kinetics and mechanisms of precipitation at 300 °C were studied by Bon ⁇ ina and Zupani ⁇ [10].
  • the invention Manufacturing of High Strength and Heat Resistant Aluminium Alloys Strengthened by Dual Precipitates is a method for manufacturing of high strength and temperature resistant Al-alloys that are strengthened by quasicrystalline and L12 precipitates. Al-alloys produced in this way are convenient, but not limited, for the application in the automobile and aerospace industries, as well as in the construction and building.
  • the required properties can be obtained by an appropriate combination of alloy's chemical composition, solidification of the melt by casting methods, in which the cooling rate in the liquid exceeds 100 K/s just before the beginning of the solidification, and two- or three stage heat treatment. Before heat treatment, the casting products can be cold or warm plastically deformed by any forming technology.
  • the invention is defined in the appended claims.
  • the heat-treatable aluminium alloy having a high strength and heat comprises Mn: 2.0-5.0 mass.%; Cr: 0.001-2.0 mass.%; Cr + Mn: 2.0-5.0 mass.%; V: 0.001-0.5 mass.%; Cu: 2.0-4.5 mass.%; Be: 0.001-0.9 mass.%; Sc: 0.05-0.5 mass.%; and comprising of at least one element out of Zr: 0.001-0.4 mass.%; Y: 0.001-0.4 mass.%; Ti: 0.001-0.4 mass.%; Hf: 0.001-0.4 mass.% and Nb: 0.001-0.4 mass.%; the balance being Al and inevitable impurities up to 0.5 mass.%.
  • the alloy with the composition given above provides Al-rich matrix containing predominantly icosahedral quasicrystalline and L12 precipitates after appropriate casting and heat treatment.
  • the alloying elements Mn, Cr and V are necessary to obtain quasicrystalline phase during solidification.
  • the quasicrystal formation is possible by either the addition of individual elements (Mn, Cr or V) or their combinations [ K. Mihama, N. Tanaka, Y. Saito: Proceedings of the China-Japan seminars Quasicrystals, Tokyo 1989, Beijing 1990, World Scientific ; D.W. Lawther, D.J. Lloyd, R.A.Dunlap: Materials Science and Engineering A, 123(1990)33-38 ].
  • the elements Cr and V can cause during heat treatment the formation of Cr- and V-rich shells around cores of the IQC-precipitates that are rich in Mn, which improve the heat resistance of the invented alloy.
  • a small amount of Be enables the formation of icosahedral quasicrystalline phase by casting when cooling rates exceed 100 K/s, and it also promotes the formation of quasicrystalline precipitates during heat treatment.
  • the elements scandium and X element group (yttrium, zirconium, titanium, hafnium, and niobium) are essential for the formation of L 12 precipitates, having an ordered crystal structure.
  • the combination of Sr with any of the X-elements provide the creation of the so-called core-shell Al 3 (Sr, X) precipitates.
  • the scandium is predominantly in the cores, while the X-elements (yttrium, zirconium, titanium, hafnium, and niobium) are in their shells.
  • the presence of Cu allows the formation of Cu-rich precipitates ( ⁇ " or ⁇ ') in the matrix, between quasicrystalline and L 12 precipitates during the third, low-temperature heat treatment. It also causes the in-situ conversion of the quasicrystalline phase to the T-phase (Al 20 Mn 3 Cu 2 ) at higher temperatures during service.
  • the alloy can be produced in the liquid state by melting in induction or other furnaces, preferably under a vacuum or protective atmosphere to prevent loss of some alloying elements, hydrogenation of the melt and formation of oxides.
  • technically pure aluminium and commercial master alloys e.g. AlCu50, AIY10 can be used.
  • Transformation to the solid state should be done by any casting technique, by which a cooling rate exceeds 100 K/s. These methods encompass but are not limited to casting into permanent moulds, gravitational casting, high-pressure die-casting, centrifugal casting, continuous casting, single-roll and dual-roll casting. Adequate cooling rates firstly provide the formation of quasicrystalline phase within grain boundaries and at the grain boundaries than prevent grain growth during further heat treatments. Secondly, a sufficient amount of alloying elements is preserved in the Al-rich solid solution, which enables precipitation of different types of precipitates during further heat treatment (a modified T5 heat treatment).
  • the alloy with the appropriate composition being solidified with the suitable cooling rate has a reasonable ductility so the castings can be plastically deformed before heat treatment by any forming process.
  • the plastic deformation can increase the alloy's homogeneity, induce fragmentation of phases formed during solidification, and can speed up the precipitation processes during ageing. This step is optional.
  • Heat treatment consists of two or three steps of artificial ageing, which represent a modified T5 heat treatment.
  • the aim of ageing at a first ageing temperature for a first predetermined time is to obtain an exceptionally high density of icosahedral quasicrystalline precipitates.
  • a quasicrystal represents a state of matter, in which atoms are regularly arranged in space, but are not periodic.
  • An icosahedral quasicrystal is quasiperiodic in three dimensions - in space. It also has a very high symmetry. Hence, it can often grow if the spherical form which advantageous in comparison to plate-like or needle-like growth of many other phases in aluminium alloys.
  • the quasiperiodicity allows excellent matching of icosahedral quasicrystal with the aluminium matrix. Therefore, the interface energy between the matrix and icosahedral precipitates is very low.
  • the low interface energy enables homogeneous nucleation of IQC-precipitates, which are formed in very high number density.
  • the size of IQC-particles can be 10-15 nm, and the interparticle distances are 30-40 nm.
  • the mobility of Mn atoms determines the size and interparticle distances between IQC-precipitates. For this reason, temperatures between 260 °C and 340 °C are required, and the ageing times are between 2 hours and 90 hours.
  • the quasicrystalline precipitates contain aluminium and manganese predominantly.
  • cuboidal Al 3 Sc-precipitates can form in between IQC-precipitates. Their sizes are below 10 nm.
  • the microstructure after first ageing treatment consists of a mixture of nanoscale quasicrystalline and L 12 precipitates in the Al-rich solid solution.
  • the goal of the second ageing is to make the microstructure more stable when exposed to higher temperatures during service.
  • the second ageing temperature should be higher than the first ageing temperature.
  • the second ageing is to be carried out in the temperature range between 350 °C and 490 °C, and the duration from 15 minutes up to 10 hours.
  • the quasicrystalline precipitates that formed during first ageing stem continue to grow, the size becomes greater, and they also coarsen.
  • the number density of IQC-precipitates decreases.
  • an alloy contains Cr and V
  • a Cr- and V-shell form around the Mn-rich core.
  • the shell retards the coarsening of the IQC-precipitates during exposure of the alloy to higher temperatures.
  • L 12 precipitates, rich in scandium (e.g. Al 3 Sc) will continue to grow.
  • the elements X Zr, Y, Ti, Hf, Nb
  • the shell rich in X-elements will make harder the coarsening of L 12 -precipitates.
  • the microstructure after the second ageing consists of the core-shell icosahedral precipitates, with the sizes 20-50 nm, and the core-shell L 12 -precipitates, with the sizes 10-20 nm, that are uniformly distributed between IQC-precipitates.
  • This microstructure is very stable up to the temperatures close to the temperature of the second ageing treatment (maximum 450 °C).
  • the third ageing treatment is optional. Its aim is to cause precipitation of Cu-rich precipitates from the Al-matrix in the spaces between the icosahedral and L 12 precipitates. It is required that the alloy is quenched from the second ageing temperature, and then heated to a third ageing temperature, which is between 120 °C and 195 °C. The third ageing time is between 2 hours and 16 hours. The content of Cu in IQC- and L 12 -precipitates is minimal. Therefore, copper remains in the solid solution at the temperature of the second ageing. By quenching, it stays in the solid solution. After heating to the third ageing temperature, it precipitates in the form of Cu-rich precipitates; ⁇ ' and ⁇ ". By this treatment, the strength of the alloy at a low temperature in increased. However, the heat resistance is not changed because the Cu-rich precipitates dissolve when the alloy is heated to a temperature above 200 °C.
  • the important advantage of this method is that different shapes of castings and the desired as-cast microstructure that allows the successful heat treatment can be obtained by typical casting methods in aluminium foundries.
  • the contents of alloying elements are rather low.
  • the cast parts can be plastically deformed, if necessary, due to their ductility.
  • the quantities of alloying elements, especially those with higher prices (e.g. scandium), are low.
  • the cost per unit of property is very low. This is important in comparison with scandium alloyed aluminium alloys (e.g. patent EP 2598664 B1 ).
  • the IQC- and L 12 -precipitates are coherent with the Al-matrix, thus can be formed in much higher number densities than dispersoids. Thus, there are many more possibilities to tune the properties of the alloy by heat treatment and plastic deformation.
  • the chemical composition of the alloy was: mass. % Mn mass. % Be mass. % Cu mass. % Zr mass. % Sc mass. % Y mass. % Si mass. % Fe mass. % Al 3.29 0.36 3.57 0.11 0.51 0.12 0.21 0.16 remain
  • the alloy was produced by melting technically pure aluminium (Al 99.5), and master alloys AlMn10, AlBe5.5, AlCu50, AlZr10, AISc2 and AlY10 in an electric resistant furnace.
  • the alloy melt was homogenized at 750 °C and cast into a copper mould having cylindrical parts with diameters of 2.5, 4, 6 and 10 mm (the hardness in the as-cast condition was 100-120 HV 1).
  • Samples were heat treated in an electrical resistance furnace in air.
  • the first ageing temperature was 300 °C, and the first ageing time was 30 hours.
  • the hardness of the cast and aged samples was 120-130 HV 1
  • the hardness of the cast, deformed and aged samples was 130-140 HV 1.
  • the samples were aged for the second time in an electrical resistance furnace in air.
  • the second ageing temperature was 400 °C and the second ageing time was 1 hour.
  • the hardness of the cast, deformed and twice aged samples was 100-110 HV 1
  • the hardness of the cast, deformed and twice aged samples was 110-120 HV 1.
  • the samples were quenched in water from the second ageing temperature. Then they were aged for the third time in an electrical resistance furnace in air.
  • the third ageing temperature was 170 °C and the third ageing time was 5 hours.
  • the hardness of the cast and three times aged samples was 120-130 HV 1, and the hardness of the cast, deformed and three times aged samples was 130-140 HV 1.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Continuous Casting (AREA)
  • Forging (AREA)
  • Powder Metallurgy (AREA)

Description

    The technical field of the invention
  • The invention belongs to the area of non-ferrous metallurgy, more specifically to the casting and heat treatment of aluminium alloys.
    • Description of the background art
  • The aluminium and aluminium alloys are characterized by low densities, high specific strengths (strength divided by density), and competitive corrosion resistance. Aluminium parts can be produced economically using versatile manufacturing technologies. By using standard strengthening mechanisms, limiting values of tensile strengths between 550 MPa and 600 MPa can be achieved in wrought alloys, but with reduced elongation (around 10 %). In Al-Si casting alloys alloyed with different alloying additions (e.g. AlSi12CuNiMg), tensile strengths up to 400 MPa can be attained, and the elongation is around 5 %.
  • Most of the high-strength aluminium alloys are precipitation hardened. The most typical heat treatment is designed by T6. The T6 heat treatment consists of solution treatment at temperatures above the alloy's solvus temperature. This treatment results in homogenising the alloy and transferring the precipitates-forming alloying elements from different intermetallic phases into an Al-rich solid solution. Alloy is then quenched. The cooling rate is to be sufficient to retain almost all dissolved alloying elements in the solid solution. During natural ageing (at room temperature) or artificial ageing (at elevated temperatures, typically below 200 °C), the alloying elements start to precipitate from the solid solution. The precipitates strengthen the aluminium alloy.
  • Heat treatment T5 is an alternative to T6. When this treatment is applied, the alloy is directly heated from temperature to the temperature of artificial ageing without solution treatment. The prerequisite for this is, for example by casting, that sufficiently high-cooling rates are achieved, which retain a high content of alloying elements in the solid solution. Precipitates are formed during direct ageing.
  • In both T5 and T6 treated alloys, precipitates strengthen the matrix at lower temperatures. At temperatures approaching 200 °C, precipitates start to coarsen and dissolve, and mechanical properties start to drop rapidly.
  • For advanced applications, aluminium alloys should be manufactured in different shapes, having an appropriate combination of low density, high-strength at room temperature, heat resistance, and competitive price. There have been several solutions in the scientific literature and patents to the above-stated problem, which have been applied in the industrial practice, and are relevant to the current topic.
  • High-strength and heat-resistant precipitation hardening alloys without Sc, loose their heat resistance at temperatures slightly above 200 °C because the precipitates become coarser or even dissolve themselves. Precipitation-hardened alloys with Sc have higher heat resistance, but are more expensive because of Sc. The content of Sc in these alloys is rather high. Because of scarcity of Sc, it is not possible to produce alloys in higher amounts. The alloys, strengthened only by quasicrystals, can be produced by rapid solidification; therefore, their use is rather limited.
    • a) High-strength and/or heat-resistant heat-treatable aluminium alloys without Sc
  • High-strength aluminium alloys are based mainly on heat-treatable alloys from the systems Al-Cu (series AA2xxx), Al-Si-Mg (series AA6xxx), Al-Li and Al-Zn-Mg (series AA7xxx). The alloys are produced by different techniques (casting, powder metallurgy, forming), and then heat-treated mainly by T6 and T5 heat treatments. Increased heat resistance is achieved by the addition of Mn, Cr, Zr, Nb, Hf or V, which can form different types of precipitates that are less soluble in the aluminium matrix, and can remain in the alloy at elevated temperatures.
  • The patent WO 2002063059 A1 discloses a two-stage ageing process of an article made from an alloy comprising at least aluminium and copper. The patent WO 2008003503 A2 discloses the method of manufacturing of AA2xxx alloys for manufacturing aluminium wrought products in relatively thick gauges. In both alloys, different Al-Cu precipitates are formed during manufacturing that cannot withstand temperatures above 200 °C.
  • The patent US 5759302 A describes the compositions of Al-Cu and Al-Zn-Mg alloys, alloyed with Mn, Cr or Zr, and heat treating procedures to obtain dispersoids from the systems Al-Mn, Al-Cr or Al-Zr that can contribute to increased fracture toughness, fatigue resistance and formability.
  • The patent US 20170051383 A1 determines the composition of the alloys and parameters of the three-step ageing process to obtain dual precipitates in the aluminium matrix that can provide the application of the said alloys for high-temperature applications. The alloy is alloyed with Cu and Zr, and with Nb, Hf or V.
  • The patent US 5226983 A discloses the Al-Zr-Li-X alloys that are produced by rapid solidification. The alloy is then subjected to multiple ageing treatments after solution treatment. The microstructure of the alloy is composed of aluminium matrix and Al3(Zr, Li) precipitates.
  • The patent WO 2011122958 A1 optimises the Al-Mg-Si-Cu alloy for high-temperature stability. It defines the chemical compositions of the alloy. The aim is to obtain an L-phase as a dominant precipitate type about its number density in the over-aged state.
  • The patent US 6074498 A defines Al-Cu-Li-Sc alloys, which are subjected to a dual ageing treatment. The purpose is to attain an array of fine T1 phase precipitates within the aluminium grain, leaving a substantially T1 phase precipitate-free zone along the grain boundaries, and an array of coarse θ' and δ' phase precipitates throughout the grains with little or no θ' and δ' phase-free zones.
    • b) High-strength and heat-resistant Al-alloys, containing Sc and L 12 precipitates
  • L12-precipitates have an ordered cubic crystal structure with a general formula Al3X, where X presents one or more elements, the most important are Sc, Zr, and Y. The L12-precipitates are coherent with Al-rich solid solution. They are formed during heat treatment, and are rather stable at elevated temperatures. Table 1 gives a list of patents for Al-alloy containing L12-precipitates. The patents disclose the composition, manufacturing procedures, e.g. casting, rapid solidification, rolling, welding, extrusion, as well as their heat treatment of aluminium alloys.
  • The patent US 6248453 B1 discloses the aluminium alloys with Sc, Er, Lu, Yb, Tm and U, and with at least one element selected from the group consisting of Mg, Ag, Zn, Li, and Cu. The alloy produced with the said composition and rapid solidification comprises of Al-matrix and Al3X precipitates with the L12-structure.
  • There are several patents with the title "High strength L12 aluminium alloys" (see Table 1). The patents disclose several basic compositions of the alloys, with the addition of at least one of scandium, erbium, thulium, ytterbium, and lutetium; and at least one of gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. One of the stated possible composition is Al-(1-8)Cu-(0.2-4)Mg-(0.1-0.5)Sc-(0.05-1.0)Zr ( US 7909947 ). In other patents, different basic alloys are proposed [in US 8002912 B2 , the base alloy is Al-(4-25)Ni-(0,1-15)Tm (transition metal), and in US 7871477 B2 , the base alloy is Al-(3-6)Mg-(0,5-3)Li]. In all cases, the microstructure after heat treatment composes of Al-rich matrix and Al3X precipitates with L12-structure.
  • The patent US 5620652 A discloses the aluminium alloys containing scandium with zirconium additions that can be used for recreational products, ground transportation, marine and aerospace structures. No details of heat treatment are given. It can be supposed that they predominantly consist of Al-rich matrix and Al3X precipitates with L12-structures.
  • The patent US 20100143185 A1 discloses a method for producing high-strength aluminium alloy powder containing L12 intermetallic dispersoids. The powder is produced by melt atomization. The patented compositions of alloys cause the formation of L12-precipitates in the powder, which are uniformly distributed in Al-rich matrix. Table 1: List of patents for Al-alloys containing L12 precipitates
    EP 2598664 B1 HIGH TEMPERATURE SCANDIUM-CONTAINING ALUMINIUM ALLOY WITH IMPROVED EXTRUDABILITY
    RU 2 446 222 HEAT-RESISTANT ALUMINIUM-BASED ALLOY AND METHOD FOR OBTAINING DEFORMED SEMI-FINISHED PRODUCTS FROM IT
    RU 2 287 600 ALUMINUM-BASE MATERIAL
    US7811395 , US7871477 , US7909947 , US8002912 , US20090260722 , US20090260723 , US20090263273 High strength L12 aluminum alloys
    US7875131 L12 strengthened amorphous aluminum alloys
    US7875133 , US20090260724 , US20090260725 , US20110041963 , US7883590 Heat treatable L12 aluminum alloys
    US7879162 High strength aluminum alloys with L12 precipitates
    US8017072 Dispersion strengthened L12 aluminum alloys
    US8409373 L12 aluminum alloys with bimodal and trimodal distribution
    US8409496 Superplastic forming high strength L12 aluminum alloys
    US8409497 Hot and cold rolling high strength L12 aluminum alloys
    US8728389 Fabrication of L12 aluminum alloy tanks and other vessels by roll forming, spin forming, and friction stir welding
    US8778098 Method for producing high strength aluminum alloy powder containing L12 intermetallic dispersoids
    US8778099 Conversion process for heat treatable L12 aluminum alloys
    US9127334 Direct forging and rolling of L12 aluminum alloys for armor applications
    US9194027 Method of forming high strength aluminum alloy parts containing L12 intermetallic dispersoids by ring rolling
    US9611522 Spray deposition of L12 aluminum alloys
    US20090263266 L12 strengthened amorphous aluminum alloys
    US20110044844 Hot compaction and extrusion of I12 aluminum alloys
    US20110052932 Fabrication of 112 aluminum alloy tanks and other vessels by roll forming, spin forming, and friction stir welding
    US20110061494 Superplastic forming high strength I12 aluminum alloys
    US20110064599 Direct extrusion of shapes with I12 aluminum alloys
    US20110085932 Method of forming high strength aluminum alloy parts containing I12 intermetallic dispersoids by ring rolling
    US20110088510 Hot and cold rolling high strength L12 aluminum alloys
    US20110091346 Forging deformation of L12 aluminum alloys
  • The patents RU 2 446 222 and RU 2 287 600 deal with deformable aluminium alloys, with scandium, zirconium, manganese, copper and vanadium. The comparison of the patented chemical composition with the chemical composition of the present application is given in Table 2. The composition ranges for copper and manganese do not overlap. Also the content of zirconium in the current application is smaller than in the Russian patents. In the current application, zirconium can be replaced by the other elements (Y, Ti, Hf, or Nb). The alloys in the patents RU 2 446 222 and RU 2 287 600 do not contain beryllium, while in this application, the necessary content of Be lies between 0.001 and 0.9 mass.% Be. Thus, in the alloys in the patents RU 2 446 222 and RU 2 287 600 there are no icosahedral quasicrystalline precipitates, since the beryllium is necessary for their formation. The patents RU 2 446 222 and RU 2 287 600 do not require a two-stage heat treatment, which is necessary in the current application, and do not demand a minimum cooling rate during solidification. Table 2: The Comparison of the chemical compositions of aluminium alloys in the patents RU 2 446 222 and RU 2 287 600 with the current application
    Element Patents RU 2 446 222 and RU 2 287 600 / mass. % This application/mass.%
    copper 1.2-2.4 2.5-4.5
    manganese 1.2-2.2 2.0-5.0
    zirconium 0.15-0.6 Zr: 0.001-0.4 mass.%; can be replaced by Y: 0.001-0.4 or Ti: 0.001-0.4 or Hf: 0.001-0.4 or Nb: 0.001-0.4
    beryllium - 0.001-0.9
    vanadium 0.01-0.15 0.001-0.5
    scandium 0.01-0.2 0.05-0.5
    chromium - 0.001-2.0
    aluminium remaining remaining
  • Belov et al. [1] developed an Al-Cu-Mn-Zr-Sc alloy, which does not require solution treatment and quenching for obtaining high-strength. This alloy contains L12 precipitates, but do not have icosahedral quasicrystalline precipitates. Instead of IQC-precipitates, Al20Mn3Cu2 precipitates formed. Their disadvantage is that they are in the form of rods. In addition, their number density is much lower than the density of icosahedral precipitates with spherical morphology.
    1. [1] N.A. Belov, A.N. Alabin, I.A. Matveeva, Optimization of phase composition of Al-Cu-Mn-Zr-Sc alloys for rolled products without requirement for solution treatment and quenching, J. Alloy. Compd., 583 (2014) 206-213.
    c) Quasicrystal-strengthened Al-alloys
  • Quasicrystals are crystals having long-range order, but without periodicity. This is the main difference with periodic crystals. Aluminium quasicrystal-strengthened alloys possess periodic aluminium-rich matrix, in which a quasicrystalline phase is dispersed. The quasicrystalline phase can form during rapid solidification and casting, or it is added to aluminium by methods typical for manufacturing composites. For the production of composites, the quasicrystalline particles can be added to an aluminium melt; parts are then obtained by casting. Alternatively, the quasicrystalline powder is added to aluminium powder, and products are produced via powder metallurgical routes.
  • The patent US 5593515 A discloses the manufacturing of alloy with a general formula Albal Qa Mb Xc Td, wherein Q represents at least one element selected from the group consisting of Mn, Cr, V, Mo and W; M represents at least one element chosen from the group consisting of Co, Ni, Cu and Fe; X represents at least one element selected from rare earth elements including Y or Mm; T represents at least one element selected from the group consisting of Ti, Zr and Hf; and a, b, c and d represent the following atomic percentages: 1≦a≦7, 0>5, 0>c≦5 and 0>d≦2. The alloy of this invention is said to be excellent in the hardness and strength at both room temperature and a high temperature, and also in thermal resistance and ductility. The products can be of a rapidly solidified alloy, a heat-treated alloy obtained by heat-treating a rapidly solidified alloy, and a compacted and consolidated alloy material obtained by compacting and consolidating a rapidly solidified alloy.
  • The patent US 5419789 A discloses the aluminium-based alloy which consists Al and 0.1 to 25 atomic % of at least two transition metal elements and has a structure in which at least quasicrystals are homogeneously dispersed in a matrix composed of Al or a supersaturated Al solid solution. The alloy is in the form of a rapidly solidified material, a heat treated material of the rapidly solidified material, or a compacted and consolidated material formed from the rapidly solidified material.
  • Schurack et al. [2, 3] wanted to achieve improved properties of aluminium alloys applying quasicrystal-strengthening. Alloys were produced by melt spinning, mechanical alloying and conventional casting. They reported that mechanical alloying of stable AlCuFe-quasicrystals and aluminium powders did not give required combination of strength and ductility. However, Ce-addition to Al-Mn-alloys improved the quasicrystal forming ability, presumably due to stabilisation of the icosahedral melt structure. This procedure allowed the direct formation of icosahedral quasicrystals while cooling the melt. The milled and extruded melt-spun ribbons of the alloy Al92Mn6Ce2 attained strength of approximately 800 MPa and elongation ∼25 %, but the conventionally cast rods had strength around 500 MPa and elongation ∼20 %.
  • Song et al. [4] found out that the addition of Be sharply reduced the critical cooling rate for formation of quasicrystals during solidification and the required Mn-content in Al-Mn alloys. They established that quasicrystals formed at conventional casting methods as well (e.g. die casting). However, in their alloys were always present additional intermetallic phases, which reduce mechanical properties. Improved version of their alloys was prepared by Rozman et al. [5].
    • [2] F. Schurack, J. Eckert, L. Schultz: Synthesis and mechanical properties of cast quasicrystal-reinforced Al-alloys, Acta materialia, 49 (2001) 1351-1361
    • [3] F. Schurack, J. Eckert, L. Schultz: Synthesis and mechanical properties of quasicrystalline Al-based composites, Quasicrystals, Structure and Physical Properties, Wiley-VCH GmbH & Co. KgaA, Weinheim, 2003, 551-569
    • [4] G.S. Song, E. Fleury, S.H. Kim, W.T. Kim, D.H. Kim: Enhancement of the quasicrystal-forming ability in Al-based alloys by Be-addition, Journal of Alloys and Compounds, 342 (2002) 251 - 255
    • [5] N. Rozman, J. Medved, F. Zupanič, Microstructural evolution in Al-Mn-Cu-(Be) alloys, Philos. Mag., 91 (2011) 4230-4246.
    d) Quasicrystalline precipitates in Al-alloys
  • Kim et al. [6] found decagonal quasicrystalline precipitates in as-rapidly solidified and annealed commercial AISI 2024 aluminium alloy containing 2 wt% Li. The decagonal quasicrystalline phase precipitated from the solid solution at 400 °C. Zhang et al. [7, 8] studied the annealing of supersaturated solid solutions in some rapidly solidified Al-transition metal alloys using transmission electron microscopy. They found that icosahedral quasicrystalline precipitates formed in the Al-matrix in the binary systems Al-Cr and Al-Fe, and also in the systems Al-V, Al-Cr and Al-Mo, when they were alloyed by a small amount of Fe.
  • Zupanic et al. [9] discovered icosahedral precipitates in an Al-Mn-Be-Cu that was cast into a copper mould, after 24 h of annealing at 300 °C. The kinetics and mechanisms of precipitation at 300 °C were studied by Bončina and Zupanič [10].
  • There also exists a patent regarding quasicrystalline precipitates ( US 5632826 ). However, it is related to the iron-based alloy and is inapplicable to the current invention.
    • [6] D.H. Kim, K. Chattopadhyay, B. Cantor, QUASI-CRYSTALLINE AND RELATED CRYSTALLINE PHASES IN A RAPIDLY SOLIDIFIED 2024-2LI ALUMINUM-ALLOY, Acta Metallurgica Et Materialia, 39 (1991) 859-875.
    • [7] X.D. Zhang, Y.J. Bi, M.H. Loretto, STRUCTURE AND STABILITY OF THE PRECIPITATES IN MELT SPUN TERNARY AL-TRANSITION-METAL ALLOYS, Acta Metallurgica Et Materialia, 41 (1993) 849-853.
    • [8] X.D. Zhang, M.H. Loretto, Stability and decomposition mechanisms of supersaturated solid solutions in rapidly solidified aluminium transition metal alloys, Materials Science and Technology, 12 (1996) 19-24.
    • [9] ZUPANIČ, Franc, WANG, Di, GSPAN, Cristian, BONČINA, Tonica. Precipitates in a quasicrystal-strengthened Al-Mn-Be-Cu alloy. Materials characterization, ISSN 1044-5803. [Print ed.], Aug. 2015, vol. 106, str. 93-99. http://www.sciencedirect.com/science/article/pii/S1044580315001606, doi: 10.1016/j.matchar.2015.05.013.
    • [10] BONČINA, Tonica, ZUPANIČ, Franc. In situ TEM study of precipitation in a quasicrystal-strengthened Al-alloy. Archives of metallurgy and materials, ISSN 1733-3490, 2017, vol. 62, iss. 1, str. 5-9. http://www.imim.pl/files/archiwum/Vol1_2017/01.pdf, doi: 10.1515/amm-2017-0001.
    Description of a new solution
  • The invention Manufacturing of High Strength and Heat Resistant Aluminium Alloys Strengthened by Dual Precipitates is a method for manufacturing of high strength and temperature resistant Al-alloys that are strengthened by quasicrystalline and L12 precipitates. Al-alloys produced in this way are convenient, but not limited, for the application in the automobile and aerospace industries, as well as in the construction and building.
  • The required properties can be obtained by an appropriate combination of alloy's chemical composition, solidification of the melt by casting methods, in which the cooling rate in the liquid exceeds 100 K/s just before the beginning of the solidification, and two- or three stage heat treatment. Before heat treatment, the casting products can be cold or warm plastically deformed by any forming technology. The invention is defined in the appended claims.
  • The heat-treatable aluminium alloy having a high strength and heat comprises
    Mn: 2.0-5.0 mass.%;
    Cr: 0.001-2.0 mass.%;
    Cr + Mn: 2.0-5.0 mass.%;
    V: 0.001-0.5 mass.%;
    Cu: 2.0-4.5 mass.%;
    Be: 0.001-0.9 mass.%;
    Sc: 0.05-0.5 mass.%; and comprising of at least one element out of
    Zr: 0.001-0.4 mass.%;
    Y: 0.001-0.4 mass.%;
    Ti: 0.001-0.4 mass.%;
    Hf: 0.001-0.4 mass.% and
    Nb: 0.001-0.4 mass.%;
    the balance being Al and inevitable impurities up to 0.5 mass.%.
  • The alloy with the composition given above provides Al-rich matrix containing predominantly icosahedral quasicrystalline and L12 precipitates after appropriate casting and heat treatment. The alloying elements Mn, Cr and V are necessary to obtain quasicrystalline phase during solidification. The quasicrystal formation is possible by either the addition of individual elements (Mn, Cr or V) or their combinations [K. Mihama, N. Tanaka, Y. Saito: Proceedings of the China-Japan seminars Quasicrystals, Tokyo 1989, Beijing 1990, World Scientific; D.W. Lawther, D.J. Lloyd, R.A.Dunlap: Materials Science and Engineering A, 123(1990)33-38]. The elements Cr and V can cause during heat treatment the formation of Cr- and V-rich shells around cores of the IQC-precipitates that are rich in Mn, which improve the heat resistance of the invented alloy. A small amount of Be enables the formation of icosahedral quasicrystalline phase by casting when cooling rates exceed 100 K/s, and it also promotes the formation of quasicrystalline precipitates during heat treatment. The elements scandium and X element group (yttrium, zirconium, titanium, hafnium, and niobium) are essential for the formation of L12 precipitates, having an ordered crystal structure. The combination of Sr with any of the X-elements provide the creation of the so-called core-shell Al3(Sr, X) precipitates. In these precipitates, the scandium is predominantly in the cores, while the X-elements (yttrium, zirconium, titanium, hafnium, and niobium) are in their shells. The presence of Cu allows the formation of Cu-rich precipitates (Θ" or Θ') in the matrix, between quasicrystalline and L12 precipitates during the third, low-temperature heat treatment. It also causes the in-situ conversion of the quasicrystalline phase to the T-phase (Al20Mn3Cu2) at higher temperatures during service.
  • The alloy can be produced in the liquid state by melting in induction or other furnaces, preferably under a vacuum or protective atmosphere to prevent loss of some alloying elements, hydrogenation of the melt and formation of oxides. As a charge, technically pure aluminium and commercial master alloys (e.g. AlCu50, AIY10) can be used.
  • Transformation to the solid state should be done by any casting technique, by which a cooling rate exceeds 100 K/s. These methods encompass but are not limited to casting into permanent moulds, gravitational casting, high-pressure die-casting, centrifugal casting, continuous casting, single-roll and dual-roll casting. Adequate cooling rates firstly provide the formation of quasicrystalline phase within grain boundaries and at the grain boundaries than prevent grain growth during further heat treatments. Secondly, a sufficient amount of alloying elements is preserved in the Al-rich solid solution, which enables precipitation of different types of precipitates during further heat treatment (a modified T5 heat treatment).
  • The alloy with the appropriate composition being solidified with the suitable cooling rate has a reasonable ductility so the castings can be plastically deformed before heat treatment by any forming process. The plastic deformation can increase the alloy's homogeneity, induce fragmentation of phases formed during solidification, and can speed up the precipitation processes during ageing. This step is optional.
  • Heat treatment consists of two or three steps of artificial ageing, which represent a modified T5 heat treatment.
  • The aim of ageing at a first ageing temperature for a first predetermined time is to obtain an exceptionally high density of icosahedral quasicrystalline precipitates. A quasicrystal represents a state of matter, in which atoms are regularly arranged in space, but are not periodic. An icosahedral quasicrystal is quasiperiodic in three dimensions - in space. It also has a very high symmetry. Hence, it can often grow if the spherical form which advantageous in comparison to plate-like or needle-like growth of many other phases in aluminium alloys. The quasiperiodicity allows excellent matching of icosahedral quasicrystal with the aluminium matrix. Therefore, the interface energy between the matrix and icosahedral precipitates is very low. The low interface energy enables homogeneous nucleation of IQC-precipitates, which are formed in very high number density. For example, the size of IQC-particles can be 10-15 nm, and the interparticle distances are 30-40 nm. The mobility of Mn atoms determines the size and interparticle distances between IQC-precipitates. For this reason, temperatures between 260 °C and 340 °C are required, and the ageing times are between 2 hours and 90 hours. The quasicrystalline precipitates contain aluminium and manganese predominantly. During the first ageing treatment, also cuboidal Al3Sc-precipitates can form in between IQC-precipitates. Their sizes are below 10 nm. The microstructure after first ageing treatment consists of a mixture of nanoscale quasicrystalline and L12 precipitates in the Al-rich solid solution.
  • The goal of the second ageing is to make the microstructure more stable when exposed to higher temperatures during service. The second ageing temperature should be higher than the first ageing temperature. The second ageing is to be carried out in the temperature range between 350 °C and 490 °C, and the duration from 15 minutes up to 10 hours. During the second ageing, the quasicrystalline precipitates that formed during first ageing stem continue to grow, the size becomes greater, and they also coarsen. During the coarsening the number density of IQC-precipitates decreases. When an alloy contains Cr and V, a Cr- and V-shell form around the Mn-rich core. Since the diffusivities of Cr- and V- are much smaller than that of manganese, the shell retards the coarsening of the IQC-precipitates during exposure of the alloy to higher temperatures. L12 precipitates, rich in scandium (e.g. Al3Sc) will continue to grow. During this stage, the elements X (Zr, Y, Ti, Hf, Nb) will form a shell around Sc-rich L12 precipitates. They have smaller diffusivities than Sc. Thus, the shell rich in X-elements will make harder the coarsening of L12-precipitates. The microstructure after the second ageing consists of the core-shell icosahedral precipitates, with the sizes 20-50 nm, and the core-shell L12-precipitates, with the sizes 10-20 nm, that are uniformly distributed between IQC-precipitates. This microstructure is very stable up to the temperatures close to the temperature of the second ageing treatment (maximum 450 °C).
  • The third ageing treatment is optional. Its aim is to cause precipitation of Cu-rich precipitates from the Al-matrix in the spaces between the icosahedral and L12 precipitates. It is required that the alloy is quenched from the second ageing temperature, and then heated to a third ageing temperature, which is between 120 °C and 195 °C. The third ageing time is between 2 hours and 16 hours. The content of Cu in IQC- and L12-precipitates is minimal. Therefore, copper remains in the solid solution at the temperature of the second ageing. By quenching, it stays in the solid solution. After heating to the third ageing temperature, it precipitates in the form of Cu-rich precipitates; θ' and θ". By this treatment, the strength of the alloy at a low temperature in increased. However, the heat resistance is not changed because the Cu-rich precipitates dissolve when the alloy is heated to a temperature above 200 °C.
  • The important advantage of this method is that different shapes of castings and the desired as-cast microstructure that allows the successful heat treatment can be obtained by typical casting methods in aluminium foundries. The contents of alloying elements are rather low. Thus, the cast parts can be plastically deformed, if necessary, due to their ductility. The quantities of alloying elements, especially those with higher prices (e.g. scandium), are low. Thus, the cost per unit of property is very low. This is important in comparison with scandium alloyed aluminium alloys (e.g. patent EP 2598664 B1 ). The IQC- and L12-precipitates are coherent with the Al-matrix, thus can be formed in much higher number densities than dispersoids. Thus, there are many more possibilities to tune the properties of the alloy by heat treatment and plastic deformation.
    • A reference example of the manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates
  • The chemical composition of the alloy was:
    mass. % Mn mass. % Be mass. % Cu mass. % Zr mass. % Sc mass. % Y mass. % Si mass. % Fe mass. % Al
    3.29 0.36 3.57 0.11 0.51 0.12 0.21 0.16 remain
  • The alloy was produced by melting technically pure aluminium (Al 99.5), and master alloys AlMn10, AlBe5.5, AlCu50, AlZr10, AISc2 and AlY10 in an electric resistant furnace. The alloy melt was homogenized at 750 °C and cast into a copper mould having cylindrical parts with diameters of 2.5, 4, 6 and 10 mm (the hardness in the as-cast condition was 100-120 HV 1).
  • Thereafter, some samples were plastically deformed at room temperature for 50 %.The hardness of deformed samples increased to 140-150 HV 1.
  • Samples were heat treated in an electrical resistance furnace in air. The first ageing temperature was 300 °C, and the first ageing time was 30 hours. After the first ageing, the hardness of the cast and aged samples was 120-130 HV 1, and the hardness of the cast, deformed and aged samples was 130-140 HV 1.
  • The samples were aged for the second time in an electrical resistance furnace in air. The second ageing temperature was 400 °C and the second ageing time was 1 hour. After the second ageing, the hardness of the cast, deformed and twice aged samples was 100-110 HV 1, and the hardness of the cast, deformed and twice aged samples was 110-120 HV 1.
  • The samples were quenched in water from the second ageing temperature. Then they were aged for the third time in an electrical resistance furnace in air. The third ageing temperature was 170 °C and the third ageing time was 5 hours. After the third ageing, the hardness of the cast and three times aged samples was 120-130 HV 1, and the hardness of the cast, deformed and three times aged samples was 130-140 HV 1.

Claims (2)

  1. Method for Manufacturing of High Strength and Heat Resistant Aluminium Alloys Strengthened by Dual Precipitates,
    in which the alloys having the following chemical composition:
    Mn: 2.0-5.0 mass.%;
    Cr: 0.001-2.0 mass.%;
    Cr + Mn: 2.0-5.0 mass.%;
    V: 0.001-0.5 mass.%;
    Cu: 2.0-4.5 mass.%;
    Be: 0.001-0.9 mass.%;
    Sc: 0.05-0.5 mass.%; and comprising of at least one element out of Zr: 0.001-0.4 mass.%;
    Y: 0.001-0.4 mass.%;
    Ti: 0.001-0.4 mass.%;
    Hf: 0.001-0.4 mass.% and
    Nb: 0.001-0.4 mass.%;
    the balance being Al and inevitable impurities up to 0.5 mass.%.
    the alloys are manufactured by solidification of the melt by casting methods, in which a cooling rate in the liquid before the beginning of solidification is higher than 100 K s-1, the alloys are optionally deformed by plastic deformation and mandatory heat treated according to a two-step-ageing treatment
    wherein the first ageing at the first ageing temperature takes place between 260 °C and 340 °C, and the first ageing time is between 2 hours and 90 hours, where the microstructure after the first ageing treatment consists of quasicrystalline and L12 precipitates in the Al-rich solid solution;
    wherein the second ageing temperature is between 350 °C and 490 °C, and the second ageing time is between 15 minutes and 10 hours; and the microstructure after the second ageing consists of stabilised, coarser icosahedral precipitates, and core-shell L12 precipitates.
  2. Method for Manufacturing of High Strength and Heat Resistant Aluminium Alloys Strengthened by Dual Precipitates of Claim 1
    wherein the alloy is quenched from the second ageing temperature in water or similar cooling medium at room temperature, and thereafter aged at a third ageing temperature for a third ageing time; where the third ageing temperature is between 120 °C and 195 °C, and the third ageing time is between 2 hours and 16 hours; where Cu-rich precipitates are formed in the Al-matrix in the spaces between icosahedral and L12 precipitates to enhance a high-strength of the alloy at room temperature.
EP17468002.5A 2017-09-13 2017-12-27 Manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates Active EP3456853B1 (en)

Applications Claiming Priority (1)

Application Number Priority Date Filing Date Title
SI201700256A SI25352A (en) 2017-09-13 2017-09-13 Production of high-strength and temperature resistant aluminum alloys fortified with double excretion

Publications (2)

Publication Number Publication Date
EP3456853A1 EP3456853A1 (en) 2019-03-20
EP3456853B1 true EP3456853B1 (en) 2020-02-19

Family

ID=61024552

Family Applications (1)

Application Number Title Priority Date Filing Date
EP17468002.5A Active EP3456853B1 (en) 2017-09-13 2017-12-27 Manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates

Country Status (2)

Country Link
EP (1) EP3456853B1 (en)
SI (1) SI25352A (en)

Families Citing this family (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
AU2018379386B2 (en) 2017-12-04 2021-09-23 Monash University High strength aluminium alloy for rapid solidification manufacturing processes
CN111206194B (en) * 2020-02-24 2022-01-28 湖南人文科技学院 Thermal mechanical treatment process for preparing aluminum alloy with high comprehensive performance
CN113684393B (en) * 2020-05-22 2022-06-17 信承瑞技术有限公司 Preparation process of high-strength high-conductivity copper-selenium alloy contact wire
CN114107764B (en) * 2020-08-26 2022-10-21 宝山钢铁股份有限公司 Jet casting and rolling 7XXX aluminum alloy thin strip and preparation method thereof
CN112410624B (en) * 2020-11-05 2021-06-15 山东迈奥晶新材料有限公司 Al-Si alloy, preparation method thereof and heat sink of 5G communication base station
RU2753537C1 (en) * 2021-02-04 2021-08-17 Общество с ограниченной ответственностью "Научно-производственный центр магнитной гидродинамики" Alloy based on aluminum for production of wire and method for obtaining it
CN113789453B (en) * 2021-08-17 2023-08-01 东南大学 Method for improving high-temperature strength of heat-resistant aluminum alloy through Mn microalloying
CN114737142B (en) * 2022-06-13 2022-08-30 中国航发北京航空材料研究院 Low-stress corrosion-resistant preparation method for aluminum alloy casting through vibratory aging and cryogenic treatment
CN117443982B (en) * 2023-11-16 2024-04-19 广州航海学院 Heat-resistant aluminum alloy wire material and preparation method thereof

Family Cites Families (19)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5226983A (en) 1985-07-08 1993-07-13 Allied-Signal Inc. High strength, ductile, low density aluminum alloys and process for making same
JP3142659B2 (en) 1992-09-11 2001-03-07 ワイケイケイ株式会社 High strength, heat resistant aluminum base alloy
SE508684C2 (en) 1993-10-07 1998-10-26 Sandvik Ab Precision-hardened iron alloy with quasi-crystalline structure particles
JPH07238336A (en) * 1994-02-25 1995-09-12 Takeshi Masumoto High strength aluminum-base alloy
JP2795611B2 (en) 1994-03-29 1998-09-10 健 増本 High strength aluminum base alloy
US5597529A (en) 1994-05-25 1997-01-28 Ashurst Technology Corporation (Ireland Limited) Aluminum-scandium alloys
JP3053352B2 (en) 1995-04-14 2000-06-19 株式会社神戸製鋼所 Heat-treated Al alloy with excellent fracture toughness, fatigue properties and formability
US6074498A (en) 1996-10-28 2000-06-13 Mcdonnell Douglas Corporation Heat treated Al-Cu-Li-Sc alloys
US6248453B1 (en) 1999-12-22 2001-06-19 United Technologies Corporation High strength aluminum alloy
US6869490B2 (en) 2000-10-20 2005-03-22 Pechiney Rolled Products, L.L.C. High strength aluminum alloy
RU2287600C1 (en) * 2005-08-09 2006-11-20 Государственное образовательное учреждение высшего профессионального образования "Московский государственный институт стали и сплавов" (технологический университет) Aluminum-base material
FR2907466B1 (en) 2006-07-07 2011-06-10 Aleris Aluminum Koblenz Gmbh ALUMINUM ALLOY PRODUCTS OF THE AA7000 SERIES AND METHOD FOR MANUFACTURING THE SAME
US7811395B2 (en) 2008-04-18 2010-10-12 United Technologies Corporation High strength L12 aluminum alloys
US8002912B2 (en) 2008-04-18 2011-08-23 United Technologies Corporation High strength L12 aluminum alloys
US7871477B2 (en) 2008-04-18 2011-01-18 United Technologies Corporation High strength L12 aluminum alloys
US8778098B2 (en) 2008-12-09 2014-07-15 United Technologies Corporation Method for producing high strength aluminum alloy powder containing L12 intermetallic dispersoids
WO2011122958A1 (en) 2010-03-30 2011-10-06 Norsk Hydro Asa High temperature stable aluminium alloy
DE102010032768A1 (en) 2010-07-29 2012-02-02 Eads Deutschland Gmbh High-temperature scandium alloyed aluminum material with improved extrudability
IN2014CH00715A (en) 2014-02-14 2015-08-21 Indian Inst Scient

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
None *

Also Published As

Publication number Publication date
EP3456853A1 (en) 2019-03-20
SI25352A (en) 2018-07-31

Similar Documents

Publication Publication Date Title
EP3456853B1 (en) Manufacturing of high strength and heat resistant aluminium alloys strengthened by dual precipitates
JP5463795B2 (en) Aluminum alloy and heat-resistant aluminum alloy material and method for producing the same
JP5239022B2 (en) High strength and high toughness magnesium alloy and method for producing the same
JP2021527758A (en) High-performance Al-Zn-Mg-Zr-based aluminum alloy for welding and additive manufacturing
KR101333915B1 (en) Aluminum-zinc-magnesium-scandium alloys and methods of fabricating same
JP3684313B2 (en) High-strength, high-toughness aluminum alloy forgings for automotive suspension parts
KR100994812B1 (en) High-strength high-ductility magnesium alloy extrudate and manufacturing method thereof
US20040191111A1 (en) Er strengthening aluminum alloy
CN113073244A (en) High-strength and high-toughness rare earth heat-resistant magnesium alloy and preparation method thereof
JP2013525608A (en) Damage-resistant aluminum material with hierarchical microstructure
US10604828B2 (en) Al—Zn alloy comprising precipitates with improved strength and elongation and method of manufacturing the same
CN109930045B (en) High-strength-toughness heat-resistant Mg-Gd alloy suitable for gravity casting and preparation method thereof
CN106929720B (en) A kind of high strength easy recrystallization wrought aluminium alloy and preparation method thereof
US20240263279A1 (en) High strength microalloyed magnesium alloy
JP2007527466A (en) Beta titanium alloy, process for producing hot rolled products from this type of alloy, and use thereof
WO2016040562A1 (en) Alloys for highly shaped aluminum products and methods of making the same
KR20160136832A (en) High strength wrought magnesium alloys and method for manufacturing the same
JP6738212B2 (en) Aluminum alloy forged product and manufacturing method thereof
US5810949A (en) Method for treating an aluminum alloy product to improve formability and surface finish characteristics
WO2021003528A1 (en) Aluminium alloys
KR101680041B1 (en) Wrought magnesium alloy having high ductility and high toughness and method for preparing the same
CN113474479A (en) Method for producing a plate or strip from an aluminium alloy and plate, strip or shaped part produced thereby
JP7459496B2 (en) Manufacturing method for aluminum alloy forgings
JP4103959B2 (en) Production method of Al-Si alloy
JPH05302138A (en) Aluminum base alloy laminated and compacted material and its manufacture

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: REQUEST FOR EXAMINATION WAS MADE

17P Request for examination filed

Effective date: 20181130

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO RS SE SI SK SM TR

AX Request for extension of the european patent

Extension state: BA ME

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: EXAMINATION IS IN PROGRESS

17Q First examination report despatched

Effective date: 20190718

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: GRANT OF PATENT IS INTENDED

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

RIC1 Information provided on ipc code assigned before grant

Ipc: C22F 1/00 20060101ALI20191106BHEP

Ipc: C22F 1/057 20060101ALI20191106BHEP

Ipc: C22C 21/00 20060101ALI20191106BHEP

Ipc: C22C 21/12 20060101ALI20191106BHEP

Ipc: C22F 1/04 20060101AFI20191106BHEP

INTG Intention to grant announced

Effective date: 20191127

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: THE PATENT HAS BEEN GRANTED

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO RS SE SI SK SM TR

REG Reference to a national code

Ref country code: CH

Ref legal event code: EP

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 602017011959

Country of ref document: DE

REG Reference to a national code

Ref country code: AT

Ref legal event code: REF

Ref document number: 1235044

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200315

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

REG Reference to a national code

Ref country code: NL

Ref legal event code: MP

Effective date: 20200219

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: NO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200519

Ref country code: RS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: FI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG4D

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: HR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: SE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: LV

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: BG

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200519

Ref country code: GR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200520

Ref country code: IS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200619

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: NL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: RO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: SM

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: ES

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: PT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200712

Ref country code: CZ

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: SK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: DK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: LT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: EE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

REG Reference to a national code

Ref country code: AT

Ref legal event code: MK05

Ref document number: 1235044

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200219

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602017011959

Country of ref document: DE

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20201120

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: AT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: IT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: PL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: SI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

REG Reference to a national code

Ref country code: DE

Ref legal event code: R119

Ref document number: 602017011959

Country of ref document: DE

REG Reference to a national code

Ref country code: CH

Ref legal event code: PL

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: MC

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

REG Reference to a national code

Ref country code: BE

Ref legal event code: MM

Effective date: 20201231

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: IE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201227

Ref country code: FR

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201231

Ref country code: LU

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201227

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: CH

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201231

Ref country code: DE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20210701

Ref country code: LI

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201231

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: TR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: MT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: CY

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: MK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

Ref country code: AL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200219

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: BE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201231

GBPC Gb: european patent ceased through non-payment of renewal fee

Effective date: 20211227

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: GB

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20211227