Characterization of phases, tensile properties, and fracture toughness in aircraft‐grade aluminum alloys

At a Glance

Metadata	Details
Publication Date	2019-05-02
Journal	Material Design & Processing Communications
Authors	Aline Emanuelle Albuquerque Chemin, Conrado Ramos Moreira Afonso, Fernando Antonio Pascoal, Carla Isabel dos Santos Maciel, Cassius Olívio Figueiredo Terra Ruchert
Institutions	Embraer (Brazil), Universidade Federal de São Carlos
Citations	20

Abstract

Aircraft-grade aluminum alloys are principal structural material for aircraft parts because they combine high strength and low weight.1, 2 Their mechanical characteristics such as high strength, toughness, and excellent anticorrosive properties stem from their microstructures that result in a combination of composition, deformation, natural and artificial aging treatments.2, 3 The 7xxx and 2xxx series aluminum alloys are hardenable by precipitation, which assigning higher mechanical properties. The microstructure in aluminum alloys has phases as intermetallic, which is usually insoluble by heat treatment. Because of it, size ranges from one to hundred microns together to the higher hardness of intermetallic than Al matrix; these phases can weaken the interface aluminum matrix—phase resulting in poor fracture toughness and strength. Dispersoids phases form during ingot homogenization it controls the grain size; precipitated particles are phases that depend on the alloying elements, their concentration, and heat treatment conditions.4 The precipitates are an obstacle to dislocations movement, which increases strength; the dislocation shear the precipitate particles if they are small or weak, or can bypass them if it is larger and stronger.5 The microstructure of AA7050 alloy promotes a properties combination as the strength, stress corrosion cracking, yielding strength and fracture toughness; hence, this alloy is used to build structure component of the wing were no have welds or rivet.5-7 Adding Zn and Mg promotes the precipitation of η’ semicoherent metastable phase. The Guinier-Preston zones I or II (GP), which are formed by natural aging at room temperature, precede η’ phase that is predominant in microstructure; the precipitation of the equilibrium η phase occurs as follows: supersaturated solid solution → vacancy rich clusters → GP zones → metastable η’ → stable η.4, 8-10 The additions of Cu on alloy forms S phase (Al2CuMg) which dissolution by multi-step solution treatment increases the corrosion resistance, the stress corrosion cracking, the strength and the damage tolerance, however, decreases hardness.4, 5, 11, 12 The Zr is added to form dispersoids Al3Zr that have similarity of aluminum crystal structure being fully coherent, form during ingot homogenized, and they act as controller recrystallization and grain growth promoting the increase of fracture toughness; however, the Al3Zr become incoherent after recrystallization.6, 13-15 The combination of the presence of phase η in the high angle grain boundary and free precipitation zone become the intergranular separation as a dominant mechanism of fracture.16 The intergranular fracture can form because of the coarse incoherent phase precipitation η in high angle grain boundary interferes the slip transfer across the grain promoting low strain in alloy under stress.5, 16 The portion of a transgranular fracture in AA7050 occur in the residual ligament of material after intergranular fracture.16 The increase of strength is successful in AA7050 by precipitation hardening and strengthening of fine grain when the alloy is solubilized and aged in the same conditions.5, 9 Both for AA2050 and AA7050 alloys the presence of Fe, Si, Cu, and Mg forms intermetallic particles that may control the fracture mechanism.5, 16 The development of AA2050 alloy resulted in low-density and high-elasticity modulus even in cryogenic conditions.1, 17 In AA2050 alloys, these phases are Al-Li (δ), Al3Li (δ’), Al2CuLi (T1), Al6CuLi2 (T2), Al15Cu8Li2 (TB), Al2CuMg (S′), Al3Zr (β’), Al2Cu (θ’), and Al2Cu (Ω).18-22 The precipitation of δ’ phase occurs during the aging of Al-Li supersaturated solid solution whose composition contains Zr and Cu.23-26 These precipitates are spherical, formed in an orderly manner in a cube-cube orientation relationship with the matrix. Because it is coherent with the Al matrix, the δ’ phase is resistant to the movement of dislocations, thus increasing the alloy’s mechanical strength.27, 28 In addition to δ’ phase, additions of Cu favor the formation of θ,’ also metastable, and Ω phases, both with Al2Cu composition, but with different crystallographic planes and morphologies.29 The θ’ phase appears as rectangular or octagonal plates on the {100} planes, while the Ω phase consists of thin hexagonal plates on the {111} planes, and are predominant in alloys with high Cu and Mg contents.19, 29 The T1 phase is considered the main phase responsible for the increase in mechanical strength of Al-Cu-Li alloys, precipitating at the dislocations and grain boundaries through the failure of the stacking sequence mechanism (crystal).19, 30 The T2 phase is an equilibrium phase that precipitates at the grain boundaries upon heat treatment at temperatures above 673 K, in alloys with high Cu: Mg ratio, nucleating predominantly in high-angle grain boundaries and, like T1, contributes to increasing the mechanical strength.31 The TB phase forms in the temperature range of 550 to 620 K; ie, if the TB phase is present, the T2 phase will not be formed.18 In the case of AA2050 alloys, the super-aging heat treatment has some peculiarities when metastable δ’ particles are involved: the long period required for this heat treatment may cause the phase to grow in size, making it no longer incoherent with the matrix, and favoring the increase in precipitate-free regions, loss of ductility, and decrease in fracture toughness.32 The key to developing alloys for specific applications is by controlling the formation, size, and distribution of phases. This paper discusses the fracture toughness at room environment, cryogenic temperature morphology of intermetallic phases, and their influence on the fractographic characteristics of AA7050-T7451 and AA2050-T84 alloys based on optical microscopy and scanning electron microscopy analyses and the morphology and distribution of phases precipitated in the microstructure of grains based on transmission electron microscopy analysis. It received the AA2050 and AA7050 alloys as 800 × 550 × 50 mm of laminated blocks heat treated, respectively, in the T84 and T7451 conditions. Table 1 describes the chemical composition. The transmission electron microscopy (TEM) was used to examine the details of the microstructure of nanometric precipitates. A plate sample was cut using a diamond disc with a thickness of approximately 500 μm and mechanically ground with 320 to 600 sandpapers to reach a depth down to 120 to 150 μm. Discs of 3 mm in diameter obtained by disc punch were sanded with 2000 sandpaper and polished with an alumina suspension (Al2O3) down to a thickness of 70 to 100 μm. Discs of 3-mm diameter obtained by disc punch were sanded with 2000 sandpaper and polished with alumina suspension (Al2O3) down to a thickness of 70 to 100 μm. Finally, the sample was submitted to dimpling down to thickness of 15 to 20 μm, following of ion milling to ensure this area for TEM analysis less than or equal to 100 nm. The intermetallic phases were analyzed using an INSPECT F50 microscope (FEI, Netherlands) to identify the images and by energy dispersive X-ray spectroscopy (EDX) (Apollo X SDD, EDAX, USA) to check their chemical composition. The sanding of the specimen with 2000 sandpaper followed by polishing with 1/4-μm diamond paste, as specified by the ASTM E3-11 standard. After polishing, the sample’s surfaces were electrolytically etched with Keller reagent to reveal the grain morphology. The tensile tests were performed as recommendations of ASTM E8-13a to determine the tensile properties on longitudinal orientation at room conditions and cryogenic temperature of 219.15 K. It calculates the yield stress (σy), using 0.2% of plastic strain, ultimate tensile stress (σU), the Young modulus (E), and cross-section reduction (%CR) and elongation (%EL). Figure 1 shows the specimens. The fracture toughness test performed at room temperature and 219.15 K, as follows ASTM E399-12. The C(T) specimen for analysis machined in TL and LT orientation as shows Figure 2. To perform the fracture toughness test in MTS servo-hydraulic testing machine, it used three specimens for each direction TL, LT. Table 2 shows the data to nucleate the crack. To perform a fracture toughness test at cryogenic temperature 219.15 K, it used a chamber with liquid nitrogen. The clip gage was used to measure the crack size using the compliance technique. Table 3 shows the ultimate tensile stress (σu), yield stress (σy), elastic modulus (E), and the elongation (%EL) for AA7050 T7451. Table 3 shows the mechanical properties for AA7050 T7451 at 296.15 and 219.15 K. The AA7050 T7451 alloy presents an increasing of 9.40% in σu, 9.86% in σy, and 9.86% in E at 219.15 K. The values of Table 3 are because of the presence of GPI (Guinier Preston zone I) precipitates, which is coherent with the Al matrix (Figure 3). These GPI zones have Zr, and A/Mg ordered as in {001}Al-fcc plane, formed starting room temperature to 150°C.. To AA7050 T7451, the yield stress is affected by solid solution strengthening from Mn and Cu even in few solid solution amounts and change by Mg that is highly soluble in Al providing an increase of yield stress.5 The deformation portion during the tensile test when alloy arrives to yield stress, the tensile stress in sequence, depends on precipitate inside the matrix of Al. The precipitate of GPI zone, which is preceded by clustering of solute atoms and vacancies, and Al3Zr when it is coherent with aluminum matrix, as shown Figure 3B, which promotes the increase of strength observed on Table 3.8, 10, 12 The nanometric precipitates size of approximately 10 nm, which is distributed uniformly on microstructure are GPI zone, Figure 3A. On Figure 3B, the η’ phase (lattice a = 0.496 nm, b = 1.403 nm) forms the microstructure of alloy AA7050 T7451 together GPI zone is mainly solid solution strengthening that increases the yield stress at 296.15 and 219.15 K.5 Table 4 shows the ultimate tensile stress (σu), yield stress (σy), elastic modulus (E), and the elongation (%EL) for AA2050 T84 alloy. The Table 4 shows stress increased at 219.15 K in 9.7% to ultimate tensile stress and 10.9% to yield stress; the elongation kept the same value for both temperatures; the elastic modulus decreased 2.67%. However, comparing the elastic modulus of these Al-Li alloy with AA7050 T7451, the AA2050 T84 presented high E value owing to the presence of δ’, Figure 4A, has a mechanism to the strengthening as modulus hardening and ordered hardening.33 Order hardening creates the antiphase boundary (APB), which extra energy to create it is a relief when the dislocations move in pairs to the second dislocation restores the disorder caused by the first.27, 33 The Young modulus of δ’ is 96 MPa but is much lower than T1 phase, that is, 175 MPa, because it hexagonal structure promotes the enhancement of strength.19, 24 Figure 4A shows the nanometric precipitates inside the grain, Figure 4B shows one of the most critical phases to increase the tensile properties on AA2050 T84 alloy. The Cu content, for example, increases the amount of T1 period, Figure 4B, that leads the increase in hardness and tensile stress.21, 24 These microstructure characteristics cause higher values of the Young modulus to AA2050 T85 than AA7050 T7451.33 Figure 5 shows the fracture toughness to LT and TL direction to AA7050 T7451 alloy at 296.15 and 219.15 K. Usually, the fracture toughness decreases with increasing yield strength, Figure 5 shows that the value for KIC is 44.7 MPa√m to LT direction, usually the fracture toughness for alloys with a percentage of recrystallized grains is at most 38.8 MPa√m that shows that AA7050 T7451 may have a minimum rate of recrystallized grains.16 The Figure 5 shows that fracture toughness to LT direction at 219.15 K is lower than 296.15 K; this behavior occurred owing because cryogenic temperature makes difficult the dislocation movement as a consequence the material can dissipate low energy. To the TL orientation, the fracture toughness is more moderate at both temperature (296.15 K and 219.15 K), because the fracture along the grain boundary above all high angle boundary is intergranular so is expected that KIC value as shows Figure 5 for TL orientation is lower.6 Figure 6 shows microstructure inside and grain boundary to AA7050 T7451 alloy. Figure 6B is a STEM micrography in a bright field which shows coarse precipitates and it heterogenic morphology. The small dot on Figure 6B is precipitate Al3Zr; the plate structure is η’ phase and precipitates on the grain boundary. Figure 6B shows the precipitate free zones which form because of the element contents decrease by precipitation process.7, 8, 34 Figure 7 shows a STEM Bright field micrograph and respective EDS elemental mapping. Figure 7 is a TEM and STEM micrograph of the AA7050-T7451 alloy that shows spherical nanometric precipitates of Al3Zr phase (with cubic unit cell crystalline structure) dispersed in the Al matrix (face-centered cubic crystal structure) in bright field (BF) images. Figure 7 shows details of grain microstructure of AA7050-T7451, such as a triple point grain boundary, precipitates inside the grains, and their chemical mapping. The phase Al3Zr is little dots that is metastable and has a spherical shape and precipitates inside the grain during ingot homogenization. The presence of Zn and Mg, Figure 7, indicates a presence of η (MgZn2) phase that precipitates along the grain boundaries, but η’ (MgZn2), Figure 3 (b). Figure 8 shows the intermetallic particles on AA7050 T7451. Figure 8 shows the morphology of the coarse intermetallics formed in the AA7050-T7451 alloys. An EDS analysis of Area 1 in Figure 8 reveals that the particle contains Al, Cu, and Fe in its composition, possibly indicating the intermetallic Al7Cu2Fe with a tetragonal crystalline structure, or Al23CuFe4, which is orthorhombic.35 These phases are harder than the matrix, do not also deform, these intermetallics “sequester” the solute that would form precipitates that improve the alloy’s mechanical properties by heat treatment. The anisotropy of microstructure of AA7050 T7451 observed in Figure 6-8 become the relationship between fracture toughness and microstructure dependent on stress state, fracture plane orientation, and temperature. Figure 9 shows the fracture toughness to AA2050 T84 at the same temperatures and plane direction as tested to AA7050 T7451 alloy. The fracture toughness for LT direction decreased 4.5% at 219.15 K whereas for TL direction the fracture toughness is independent of temperature. The adding Li to Al produces fine precipitates, which improve the stiffness and strength of the alloys, however, decreases the fracture toughness if compared with 7xxx alloys.27 Figure 9 shows the fracture toughness is lower to AA2050 T84 than AA7050 T7451 to LT direction (Figure 5). This behavior observed on both alloys is because of sizeable free precipitate zones, which decrease KIC value. However, the data at the cryogenic temperature in Table 5 shows Al-Li as a better performance alloy. The 4.5% KIC reduction for AA2050 T84 at 219.15 K is because of coherence and surface hardening cause by δ’ phase inside the grain.27 Figure 10 shows the grain microstructure of AA2050 T84 alloy. Figure 10B shows the S′ and δ’ phase distributed inside the grain, also free precipitates zones on grain boundary. The δ’ phase is resistant to the movement of dislocation, though they are shearable and cut by deformation, which decreases the effectiveness as further slip obstacle causing a coarse planar slip.32 The S′ phase has an orthorhombic crystalline structure with a = 0.405 nm, b = 0.905 nm, and c = 0.704 nm, which grows in the form of rods or thin discs along with the Al matrix in the {100} directions. For AA2050 T84 alloy composition of Cu and Mg (Cu: Mg ratio less than 3:1), the S′ is the main phase on the alloy. GP zones precede the formation of S′ phase, and Cu and Mg-rich clusters may develop within the S′ phase, or this phase may nucleate at the grain and subgrain boundaries. The strengthening precipitates shearable can results in a pileup of dislocation along the grain boundary triple junctions as shows the Figure 10A, which promotes strain localization that nucleates premature crack and results in lower fracture toughness.33 The fracture toughness can increase if strengthening precipitates dislocation looping or bypassing the particles. The amount, distribution, and morphology of precipitates and second phase particles in microstructure are significant to fracture toughness alloy, Figure 10. Figure 11 shows the chemical mapping of the AA2050-T84 alloy. Figure 11 shows a grain boundary precipitates rich in Mg, Zn, Cu, and Mn. Addition of Mg and Zn in minor amount promotes the precipitation ternary precipitates as T1 and S′/S.36 In the grain boundary, the T2 and δ phase can be precipitate.22 The larger incoherence precipitates in the grain boundary promote a loss of fracture toughness, another theory is the dislocations in those dislocations in narrow planar slip bands pile-up at grain boundaries producing high-stress concentrations at grain boundaries promoting low-energy fracture by localized shear.22, 32, 37 Figure 12 shows the intermetallic particles on the microstructure of AA2050 T84. Figure 12 shows the intermetallic phase found in the AA2050-T84 alloy, which EDS detected Mg, Cu, and Fe in the microstructure. The intermetallic phases such as those shown in Figure 12 are harder than the aluminum matrix. Because they are coarser, their size is in the order of microns; they reduce the toughness of the material, making it susceptible to fatigue crack nucleation.35 The research of commercial aircraft-grade aluminum alloys is essential to find new observation about the relationship between mechanical properties under exceptional conditions as cryogenic temperature and microstructure. The AA2050 T84 presented higher values for elastic modulus because of the very fine coherent microstructure, which is alloy strengthening. The fracture toughness of AA2050 T84 is lower than AA7050 T7451, but, at 219.15 K, the Al-Li presented higher performance, which shown that AA2050 T84 is proper to an application at low temperature. The coarse precipitate observed on the microstructure of AA7050 T7451, distributed along the grain boundary and inside the grain associated to few precipitate free zones promotes superior tensile properties as ultimate stress and yield stress at environmental temperature and cryogenic temperature compare with AA2050 T84. Dr. Aline Emanuelle Albuquerque Chemin (Chemin A.1): The main author was responsible by research of AA2050 Alloy microstructure as part of her PhD, purposed the conception of this paper and write about the formation of phases of AA 2050 alloy, showed in this paper. Dr. Conrado Moreira Afonso (Afonso C.R.M2): He is an assistant professor that helped us to investigate the phases presented in this paper by transmission electron microscopy. Dr. Fernando Pascoal (Pascoal F.3) and Dr. Carla Maciel (Maciel C.I.S.1): It were responsible to AA7050 alloy analysis of it phases on this paper. Dr. Cassius (Ruchert C.O.F.T.1): He is an assistant professor that is the head of the AA2050 research and helped to write section of this alloy. Dr. Waldek Wladimir Bose Filho (Bose Filho W.W.1): He is a professor that is the head the AA7050 research and helped to write section of this alloy and reviewed the paper to submitting 3. The authors declare that they have not conflict of interest. Ministério da Ciência, Tecnologia e Inovação > Conselho Nacional de Desenvolvimento Científico e Tecnológico.

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Original Source

DOI: https://doi.org/10.1002/mdp2.79

References

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