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Microstructure and texture development of 7075 alloy during homogenisation Abhishek Ghosh & Manojit Ghosh To cite this article: Abhishek Ghosh & Manojit Ghosh (2018): Microstructure and texture development of 7075 alloy during homogenisation, Philosophical Magazine, DOI: 10.1080/14786435.2018.1439596 To link to this article: https://doi.org/10.1080/14786435.2018.1439596

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Philosophical Magazine, 2018 https://doi.org/10.1080/14786435.2018.1439596

Microstructure and texture development of 7075 alloy during homogenisation Abhishek Ghosh and Manojit Ghosh Department of Metallurgy and Materials Engineering, Indian Institutes of Engineering Science and Technology, Howrah, India

ABSTRACT

The microstructure evolution of Al–Zn–Mg–Cu alloy during homogenisation was studied by optical microscope, field emission scanning electron microscope, energy dispersive X-ray Spectroscopy, differential scanning calorimetry and X-ray diffraction in detailed. It has been found that primary cast structure consisted of primary α (Al), lamellar eutectic structure η Mg(Zn, Cu, Al)2 and a small amount of θ (Al2Cu) phase. A transformation of primary eutectic phase from η Mg(Zn, Cu, Al)2 to S (Al2CuMg) was observed after 6 h of homogenisation treatment. The volume fraction of dendrite network structure and intermetallic phase was decreased with increase in holding time and finally disappeared after 96 h of homogenisation, which is consistent with the results of homogenisation kinetic analysis. Crystallographic texture of this alloy after casting and 96 h of homogenisation was also studied. It was found that casting process led the development of strong Goss, Brass, P and CuT components, while after homogenisation Cube, S and Copper components became predominant. Mechanical tests revealed higher hardness, yield strength and tensile strength for cast materials compared to homogenised alloys due to the presence of coarse microsegregation of MgZn2 phase. The significant improvement of ductility was observed after 96-h homogenisation, which was attributed to dissolution of second phase particles and grain coarsening. Fracture surfaces of the cast samples indicated the presence of shrinkage porosity and consequently failure occurred in the interdendritic regions or grain boundaries with brittle mode, while homogenised alloys failed under ductile mode as evident by the presence of fine dimple surfaces.

ARTICLE HISTORY

Received 3 June 2017 Accepted 5 February 2018 KEYWORDS

Al–Zn–Mg–Cu alloy; intermetallic phases; homogenisation; microstructure; crystallographic texture

1. Introduction 7075 alloys are widely used in aircraft and automobile industry due to their high strength, fracture toughness, stress corrosion cracking properties and low density

CONTACT Abhishek Ghosh

[email protected]

© 2018 Informa UK Limited, trading as Taylor & Francis Group

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A. GHOSH AND M. GHOSH

[1–4]. Owing to growing demands for achieving better mechanical properties, high strength aluminium alloys are always under focus. One of the popular or efficient ways of strengthening Aluminium alloys is primary (T4, T6) or duplex ageing treatment. For 7075 alloys, the main strength contribution comes from the nanoscale precipitates of metastable Zn and Mg rich η′ (MgZn2) phase and its precursors [5–9]. High strength alloys are used for structural applications requiring good formability [10]. However, as strength and ductility are inversely related, enhancement of strength of a material maintaining reasonable ductility is still a great challenge for the researchers. High strength 7xxx alloy can form micro-segregation and coarse constituent particles at the grain boundary during casting [11,12]. During casting, mechanical properties of 7xxx series alloys are seriously affected by the presence of different types of intermetallic phases [13]. It is well known that below solidus temperature, several major intermetallic phases namely η (MgZn2), T (Al2Mg3Zn3), S (Al2CuMg), θ (Al2Cu), (Al7Cu2Fe), are formed during solidification in Al–Zn–Mg–Cu alloys [2,11,14–16]. Subsequently, any other alloying additions may lead to the formation of more number of eutectic phases. In cast alloys, Cu, Mg, Zn solute atoms are heterogeneously distributed within the grain boundaries creating a concentration profile from the grain boundaries to grain interior, popularly known as cored structure [17,18]. Existing literature reported that η (MgZn2), (Al7Cu2Fe) phase are present in the cast structure and η is transformed to S after homogenisation in Al–Zn–Mg–Cu alloys [5,19,20]. The chemical composition of the alloy, mode of the casting process and homogenisation treatment govern the nature of coarse intermetallic phases [14]. Again, the larger volume fraction of coarse residual phases is acted as the sites of stress concentration or crack initiation and deteriorates the alloy in terms of toughness, fatigue strength and corrosion resistance [21–26]. In order to get uniform microstructure and property or otherwise remove the dendrite segregation and dissolution of coarse residual constituents, cast materials need homogenisation. Therefore, the main advantages of homogenisation process are (a) elimination of dendritic structure with reduction of volume fractions of eutectic phase (b) increment of the formability of material. Selection of homogenisation is alloy specific and must be below the solidus point. Because homogenisation is a diffusion-controlled process, the extent of diffusion is controlled by the temperature and time of the homogenisation treatment [27,28]. Changes in microstructure during this process have been assisted by changes in phase and crystallographic texture which contributes considerably in controlling the final properties of the material. Homogenisation may also be accompanied by recrystallisation (formation of new strain free grains) and grain growth, by migration of high angle grain boundaries resulting in a change of grain orientation and be responsible for development or alteration of crystallographic texture [29]. On the other hand, when the phase transformation occurred in a material, it is always associated with a change in crystal orientation relationship between the parent and the product phase. Additionally, in casting

PHILOSOPHICAL MAGAZINE

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processes or fluid-to-solid transitions generation of textured solids is unavoidable when there are enough time and activation energy for atoms to find places in existing crystals. Furthermore, in 7xxx series alloys, the grain sizes in the cast and the homogenised conditions and different types of intermetallic phases can change crystallographic texture [30]. The basic understanding of high temperature homogenisation (recovery and recrystallisation) could be understood with the information from bulk texture analysis [31–33]. The homogenisation process is known to dissolve the secondary phases and eliminate the composition segregation which improves the plasticity and subsequent processing of the deformation of as cast 7075 alloy. Therefore, it is important to examine the evolution of eutectic phases and develop the appropriate homogenisation process for industrial perspective. The objective of the paper, therefore, is to investigate the change in microstructure, texture and mechanical properties for 7075 alloy before and after homogenisation. The evolution of intermetallic phase during heat treatment is also investigated using various characterising tools. Enhancement of ductility in the cast and homogenised conditions was measured by the tensile test and subsequently, the nature of fractography was reported. 2. Materials and methods The material for the present investigation was developed by chill casting method. The chemical composition of this alloy was given in Table 1. Homogenisation treatment was carried out at 465 °C for different time durations (0 to 96 h) then cooled in air. Temperature accuracy of the furnace was maintained by a PID controller within 465 ± 1 °C. Differential scanning calorimetry (DSC) was performed with a PerkinElmer made Pyris Diamond TG/DTA (heat flow with endothermic reaction pointing down in Y-axis) and METTLER 4000-TC 11 with Perkin Elmer instrument (heat flow with endothermic reaction pointing up in Y-axis) with a constant heating rate at 10 °C/min. Samples used in DSC were mechanically ground into thin slice weighing about 10 mg. High purity copper pans were used as a reference in this study. Microstructural characterisation of the samples was performed by optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive X-ray Spectroscopy (EDS). Samples for OM and SEM were prepared by following standard metallographic polishing techniques. SEM, in the backscattered electron (BSE) mode, was performed by JEOL JSM-7600F FESEM, operated at 15 kV. EDS was performed to analyse the chemical composition of the intermetallic phases present in as cast and homogenised samples. X-ray diffraction (XRD) was conducted to identify the phases in as cast and different homogenised Table 1. Compositions of 7075 alloy (wt. %). Zn 5.38

Mg 2.44

Cu 1.59

Cr 0.18

Fe 0.10

Si 0.1

Mn 0.023

Al Bal.

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materials. It was carried out by PW 1830 generator, with scan rate 0.02°/sec using Co Kα radiations, operated at 35 kV/25 mA. The tensile tests were carried out at room temperature with a Universal Instron machine at a constant strain rate of 10−3s−1. The test samples were prepared according to ASTM E-8M standard having a gauge length of 25 mm. At least three specimens were tested and the average values were reported. The fracture surface of tensile tested samples was examined by SEM. Bulk texture measurement was carried out using a Bruker D8 Discover texture goniometer with CuKα radiation. Four incomplete pole figures (1 1 1), (2 0 0), (2 2 0), (3 1 1) on the Normal direction (ND) - transverse direction (TD) plane, keeping casting direction (CD) mutually orthogonal to ND and TD for each sample were measured. The orientation distribution function (ODF) was measured from corrected pole figure data, by the formulation of Bunge. The ODFs were represented as the plots of constant φ2 sections in Euler space defined by three Euler angles φ1, Φ, φ2.Volume fractions of main texture components were evaluated from ODF, considering an angular spread of 11° of ideal orientations in Euler space. Orthorhombic symmetry has been imposed for representation the texture for as cast and homogenised alloy. 3. Results and discussion 3.1. As cast microstructure

Figure 1(a) showed the typical as cast microstructure exhibiting serious dendritic segregation, caused by non-equilibrium solidification during casting. The average grain size of cast material is 65 μm with a standard deviation of ±5 μm. Figure 1(b) illustrated that these eutectic phases were lamellar in nature. It can also be seen that a large amount of low melting-point eutectic phases, mainly of two varieties, is present in as cast condition (Figure 1(b)). EDS result revealed the chemical compositions of white eutectic structure phase in (Figure 1(b)) bearing composition (in at%) Al-39.95, Mg-23.1, Zn-18.92, Cu-18.03 which was close to Mg(Zn, Cu, Al)2 (Figure 1(c)) [10,12,34]. The composition of other (smaller portion) gray eutectic phase in (Figure 1(b)) demonstrated the composition range (in at%): Al-57.34, Cu-38.25 which was close to the stoichiometry of Al2Cu (θ phase) (Figure 1(d)). The Mg(Zn, Cu, Al)2 phase has same crystallographic lattice as (MgZn2) [34]. So it may, be concluded that the dark areas were a primary solid solution of α (Al) with bright portions of eutectic Mg(Zn, Cu, Al)2 (η phase) and gray Al2Cu (θ phase). The most commonly observed phases in 7xxx alloy during casting were MgZn2, T (Al2Mg3Zn3). Although, a small amount of Fe affluent phases like Al7Cu2Fe and Mg2Si might also be present, depending upon the mode of casting, solidification rates and chemical compositions [2]. Figure 2 showed the main chemical elements distribution in as cast alloy. The elements were mainly populated at the vicinity of the grain boundary with a decreasing sequence from Cu, Zn to Mg which led to the formation of the concentration gradient in the


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Figure 1. (colour online) Microstructures of as cast material (a) in low magnification (b) in high magnification and corresponding their EDS results of the intermetallic phases (c) Mg (Zn, Cu, Al)2 (d) Al2Cu.

Figure 2. (colour online) SEM microstructure of as cast material in grain boundary region (high magnification) (a) and main chemical elements distribution of (b) Cu (c) Zn (d) Mg.

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alloy [15]. Therefore, the homogenisation treatment is required to remove severe dendritic segregation in cast material. The relationship between diffusion coefficient and temperature can be described by the following equation [27], D = D0 exp −

(

Q RT

)

(1)

where D0 is the diffusion co efficient, Q is the diffusion activation energy, R the gas constant and T is the absolute temperature. Temperature has a most intense influence on diffusion rate and coefficient. For higher temperature, the larger diffusion coefficients make easier elimination of dendritic segregation in cast alloy. However, the higher temperature of the homogenisation treatment increases the possibilities of over-burning of the alloy. 3.2. Microstructure of different homogenised conditions by optical microscope

The chronological changes in microstructure during homogenisation were revealed by optical microscopy (Figure 3). Large non-equilibrium constituents were still present along grain boundaries after 15 min of homogenisation treatment (Figure 3(a)). After 6 h of homogenisation, the dissolution of eutectic phase into the matrix (Figure 3(b)) was noticed. With the progress of homogenisation, the grain boundary becomes thinner and dendrite network decreased and residual phases were more sparsely distributed within the matrix (Figure 3(c)–(e)). The net-shaped eutectic structures were almost fully dissolved after 96-h homogenisation, only a few dark black phases were left (Figure 3(f)). The area fractions of the coarse particles were decreased with homogenisation times (Figure 4) [11,20]. 3.3. Evolution of intermetallic phases during homogenisation

Figure 5 represented BSE images for 7075 alloy after different homogenised times at 465 °C. The primary intermetallic phase (η) was reported to be formed by dissolving some amount of Al and Cu in MgZn2. [2,14]. After 6 h of homogenisation (Figure 5(a)–(c)) a new intermetallic phase (white in color) bearing composition (in at %) Al-47.48, Mg-22.27, Zn-4.55, Cu-25.07 with a stoichiometry close to Al2CuMg (S phase) was formed and this was noticeably absent in the cast material. The high magnification SEM image of S phase was given in (Figure 5(c)).The lamellar structure of η Mg (Zn, Cu, Al)2 obtained during casting transformed to S (Al2CuMg) after 6 h of homogenisation treatment [14]. The S phase nucleates and forms along the grain boundary between the primary η phase and the matrix [14,35]. The segregation of solutes distributed heterogeneously with a high concentration of Mg, Zn and Cu atoms inside the dendritic regions. During


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Figure 3. Optical micrograph of homogenised alloy at 465 °C for (a) 15 min (b) 6 h (c) 12 h (d) 24 h (e) 48 h and (f) 96 h.

homogenisation treatment, these solute atoms diffused from the eutectic structure of primary η phase getting it gradually dissolved into the matrix. The diffusion rate of Cu is lower than those of Zn and Mg resulted in a higher concentration of Cu in this region. The concentration of Cu at the grain boundary is enough for nucleation of S phase, making the diffusion process controlled by Cu atoms after 6 h of homogenisation [11,36]. From Figure 5(b) it can also be seen that S phase with much lower Zn content compared to primary eutectic phase, indicating faster diffusion of Zn atoms from the primary phase during homogenisation and making S is a stable phase [18,20]. After formation of S phase, it gradually grew

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Figure 4. Area percentage of coarse particles with homogenisation time at 465 °C.

and coarsened and finally took elliptical or round morphology. This occurred due to ripening and subsequent coalescence of S phase (Figure 5(d)). After 24 h of homogenisation, S phase gradually diminished making the alloy more uniform in terms of composition [14]. It can be demonstrated that by increasing holding time (to 96 h) the eutectic phase almost disappeared (Figure 5(e) and (f)) and the coarse particles become extremely fine and homogeneously dispersed throughout the alloy. The chemical composition of the coarse particles (in at %) Al-91.21, Mg-2.96, Zn-3.08, Cu-2.75 (Figure 5(f)), suggested that chemical elements uniformly distributed in the matrix. 3.4. DSC analysis of as cast and homogenised materials

Figure 6 exhibited the DSC curve of the as cast material. The two endothermic peaks on the curve at 475 and 622 °C indicated the melting point of the eutectic phase Mg(Zn, Cu, Al)2 and that of this alloy, respectively, which corroborates existing literature [11,20,34]. This has fixed the upper bound of the homogenised temperature to be 475 °C in order to eliminate over burning phenomenon of the alloy. The endothermic peak at 493 °C (Figure 7(a)) denoted the new eutectic phase S (Al2CuMg), formed after 6 h of homogenisation (enlarged view of DSC curve around S phase in (Figure 7(b)), which was further confirmed by XRD analysis. Noticeably, no peak for S (Al2CuMg) is observed in as cast alloy. The weaker intensity of the peak at 493 °C confirmed the reduction of eutectic S phase [11,20]. After 6-h homogenisation, though, the endothermic peak at 475 °C, as shown in (Figure 6(a)) of as cast material disappeared, and demonstrated the dissolution of non-equilibrium (MgZn2) phase into the matrix. The absence of any endothermic peak after 96 h of homogenisation indicated complete dissolution of phases.


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Figure 5. (colour online) Backscattered electron micrograph of the homogenised alloy at 465 °C for (a) 6 h (b) EDS results on S phase (c) enlarge view of the area containing S phase (d) 24 h (e)–(f) 96-h sample with EDS.

Therefore, it might be assumed that all non-equilibrium eutectic phases present in cast condition were dissolved fully into matrix after 96-h homogenisation. 3.5. XRD analysis

XRD results for phase evolution of as cast and different homogenised conditions (0–96 h) were shown in (Figure 8). Only α (Al) and (MgZn2) peaks appeared for as cast alloy. The small amount of Al2Cu particles untraced by XRD has been

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Figure 6. DSC curves for 7075 alloy in as cast condition.

Figure 7. (colour online) DSC curves of homogenised samples performed at 465 °C for (a) 6 h and 96 h (b) enlarged view of DSC curve of 6 h.

identified by EDS only. Conversion of η (MgZn2) to S (Al2CuMg) phase after 6 h of treatment was noticeable and in agreement with the work of Deng et al. [34]. Other peaks (e.g. η (MgZn2) and S (Al2CuMg)) also diminished gradually with the advent of homogenisation leaving only the peaks of α (Al) [14]. 3.6. Homogenisation kinetics analysis

The distribution of main chemical elements along interdendritic regions varies periodically. Therefore, the studies of diffusion law along interdendritic region are necessary to investigate the elemental distribution during homogenisation. According to diffusion kinetics for aluminium alloys during homogenisation


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Figure 8. XRD patterns of as cast and different homogenised conditions.

treatment [17,37], the initial concentration of alloying elements along dendritic region can be represented by, Fourier series components in a cosine function C(x) = C̄ + A0 cos

2𝜋x L

(2)

where, C̄ is the average concentration of the element, L is the interdendritic spacing and A0 is the initial amplitude of the composition segregation. A0 can be defined by the following equation: A0 =

) 1 1( Cmax − Cmin = ΔC0 2 2

(3)

According to Fick’s second law and the boundary conditions, A(t) is given as [12,17], ) ( 4𝜋 2 Dt A(t) = A0 exp − 2 L

(4)

By substituting Equation (1) into Equation (4), the equation can be written by, ( ))] ( 4𝜋 2 D0 t Q A(t) = A0 exp − exp − RT L2 [

(5)

Assuming the element distribution is homogenous when the composition segregation amplitude is reduced to 1%, i.e.

12


A(t) 1 = A0 100

(6)

))] ( ( 4𝜋 2 D0 t Q 1 = A0 exp − exp − 2 RT 100 L

(7)

Then, [

By taking natural logarithms of both sides, Equation (7) can be written as, ( 2 ) 4𝜋 D0 t 1 R = ln T Q 4.6L2

Assuming, A =

R , Q

B=

(

4.6 4𝜋 2 D0

(8)

) , 1 = A ln T

(

t BL2

)

(9)

Equation (9) Is the homogenisation kinetics equation. If the parameters of as cast microstructure were given, then homogenisation curves can be obtained. From the result section of 3.1, it was deduced that the diffusion coefficient of Cu is much lower than Zn and Mg at the same temperature (DCu < DMg < DZn). So homogenisation process is to be controlled by Cu [12]. By substituting the values of D0 (Cu) = 0.084 cm2/s, Q (Cu) = 136.8 kJ/mol and R = 8.314 J/mol. K into Equation (9), the homogenisation kinetic curves of Al–Zn–Cu–Mg alloy for different interdendritic spacings can be obtained as shown in Figure 9. It indicated that the soaking time was reduced dramatically with increasing the homogenisation temperature. In this work, the average interdendritic spacing L of as cast alloy was 72 μm, which was obtained from the quantitative metallographic analysis. According to the homogenisation kinetic curves, at the optimised temperatures of 465 °C, the corresponding soaking time was 96.4 h which is in good agreement with the experimental results. 3.7. Texture analysis

Macro-textures of the investigated samples were examined by high-resolution texture goniometer, based on X-ray diffractions technique. The (1 1 1) pole figures of as cast and homogenised sample (96 h) were shown in (Figure 10). The pole figures were typical of recrystallised fcc materials. It was noticeable from pole figures that homogenised samples have strong cube texture {0 0 1} . The Cast and homogenised samples can be characterised by the development of α- ( ║ND), β-fibre, γ-fibre ( ║RD) and τ-fibre [29]. Figure 11(a) and(b)


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Figure 9. (colour online) Curves showing homogenisation kinetics.

Figure 10. (colour online) {1 1 1} pole figures of (a) as cast (b) 96-h homogenised sample.

also revealed that cast and homogenised materials exhibited γ-fibre which was characterised by E {1 1 1} and F {1 1 1} texture components. The overall texture strength (T) of two samples can be found out from the following equation [38]: [ ( )]2 T = f g dg ∫ (10) Euler space where f (g) was the ODF intensity at g and g is the orientations defined by Euler angle. It was found that the intensity of ODF in cast condition, [f(g)max ] was 1.82 and the intensity of ODF in homogenised condition, [f(g)max ] is 2.02, meaning an increase of intensity of overall texture after homogenisation treatment. The volume fractions of different texture components were presented in the histogram (Figure 12). In case of cast material, strong Goss {0 1 1} , Brass {0 1 1}, P {0 1 1} and CuT {5 5 2} were the main components. On the other hand, S {1 2 3} and rotated Goss {1 1 0} were the main dominant texture components for homogenised samples. It was also noticeable that homogenisation heat treatment could also develop strong cube texture {0 0 1} due to recrystallisation of grains. For detailed understanding of texture evolution, in terms of quantitative estimation, the orientation densities of α, β, γ, τ-fibre has been plotted (Figure 13(a)– (d)). No significant change in f(g) value at Goss {0 1 1} in φ1 = 0° for as cast and homogenised samples can be noticed on α-fibre. In addition, Brass component {0 1 1} in φ1 = 35°, and rotated Goss texture component {1 1 0} in φ1 = 90° found increased slightly for as cast and homogenised samples, respectively (Figure 13(a)). The β-fibre, showed the intensity variation of Copper, S and Brass components across the φ2 section in Euler orientation. It was interesting to note that, there was no change in Copper {1 1 2} and S {1 2 3} position for both materials but at Brass {1 1 0} position the intensity value of f(g) has increased for cast sample with respect to homogenised sample (Figure 13(b)).The γ-fibres ( ‖ ND) were characterised by the E {1 1 1} and F orientations {1 1 1} in Euler space. From γ-fibre, it can be concluded that homogenised materials possess higher intensity at E and F locations compared to cast material (Figure 13(c)). The intensity variation of

Figure 13. (colour online) Texture fibre plots for (a) α-fibre (b) β-fibre (c) γ-fibre (d) τ-fibre.

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the τ-fibre was depicted in (Figure 13(d)). It was observed that, the intensity f(g) value of Goss {0 0 1} , Copper {1 1 2} , F{1 1 1} positions has been increased remarkably for homogenised material compared to cast sample, whereas the intensity of rotated cube {001} for cast material has increased slightly with respect to homogenised sample. Evolution of brass, CuT and S texture components were present due to the influence of temperature and stress fields affecting local deformation on the grains and distorted high strain lattice during casting. The P type texture, which is typical recrystallisation orientation evolved due to particle stimulated nucleation (PSN) and recrystallisation of grains during solidification and cooling [39]. In cast condition, large secondary phase particles (having size > 1 μm) (Figure 1(b)) can promote PSN by providing additional nucleation sites decreasing the recrystallised grain size. However, small dispersoids (