Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during homogenisation

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Original Article

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Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during homogenisation

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Abhishek Ghosh a,b,∗ , Manojit Ghosh a,b,∗ , Asiful H. Seikh a,b,∗

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Q1 Q2

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Q3 a Department of Metallurgy and Materials Engineering, Indian Institute of Engineering Science and Technology, Howrah 711103, India

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b

Centre of Excellence for Research in Engineering Materials, King Saud University, P.O. Box - 800, Riyadh 11421, Saudi Arabia

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a r t i c l e

i n f o

a b s t r a c t

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Article history:

Optical microscope (OM), field emission scanning electron microscope (SEM), energy disper-

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Received 16 May 2019

sive X-ray Spectroscopy (EDS), transmission electron microscope (TEM) differential scanning

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Accepted 31 August 2019

calorimetry (DSC) and X-ray diffraction (XRD) were employed to investigate the evolution of

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Available online xxx

intermetallic phases during homogenisation at 465 ◦ C for 0–48 h for Al-Zn-Mg-Cu-Sn alloy.

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Evidentially, casting is accompanied by severe dendritic segregation. The primary phases

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Keywords:

appeared during casting consisted of ␣ (Al), eutectic mixture of ␩ (Mg (Zn, Cu, Al)2 ), ␪ (Al2 Cu)

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Homogenisation

and coarse Al7 Cu2 Fe particles. After 6 h of homogenisation, two eutectic phases namely S

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Intermetallic phases

(Al2 CuMg) and Cu3 Sn were formed. The appearance of dendrites and their consequent dis-

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Dispersoids

solution with the progress of homogenisation time was modeled using a kinetic equation

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Zener pinning pressure

taking interdendritic spacing and temperature as two variants. TEM micrographs revealed

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Crystallographic texture

the presence of high density of fine dispersoids after 6 h of homogenisation following almost complete dissolution of the dendrites into the matrix after 48 h of homogenisation. Coarsening of the said dispersoids was observed with further homogenisation treatment following Ostwald ripening mechanism and consequently lowering of Zener pinning pressure (ZPP). Crystallographic texture analysis revealed the formation of strong ␥ fibre in both cast and homogenised samples. However, strong ␣ fibre along with Goss, Brass, P, and Copper components was also noticed in the cast sample. © 2019 Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

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Introduction

In recent years, the 7xxx series aluminium (Al) alloys are being widely used in automobile, aircraft industries, and space applications due to their attractive properties such as high

strength to low weight ratio, adequate toughness and stress corrosion cracking (SCC) resistance [1–3]. The most acceptable precipitation mechanism in 7xxx can be described by chronological appearance of supersaturate solid solution (SSS) - GP zone - metastable ␩’ - stable ␩ (MgZn2 ) [4]. The main strengthening phase ␩’ (fine precipitates of Zn and Mg-rich metastable

∗

Corresponding authors. E-mails: [email protected] (A. Ghosh), [email protected], manojit [email protected] (M. Ghosh), [email protected] (A.H. Seikh). https://doi.org/10.1016/j.jmrt.2019.08.055 2238-7854/© 2019 Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/ licenses/by-nc-nd/4.0/). Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055

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MgZn2 phase) is coherent with the Al matrix. The present research trend explores to extend the strength of the alloy 32 along with realistic ductility and toughness of by the addi33 tion of a small amount of trace elements like Zr, Sc, Ag, rare 34 earth, Misch metals [5,6]. However, the costs of these elements 35 remained the main bottleneck for their frequent industrial 36 use. In order to improve the age hardenability response of Al37 Cu alloy Sankaran [7], in the year 1974, probably for the first 38 time, reported the beneficial effect of Sn addition as a trace 39 element. Recently, researchers reported [8,9] that minor addi40 tion of Sn can refine the grain structure and also improve the 41 mechanical properties by forming sparsely distributed pre42 cipitates at the vicinity of the grain boundary area and held 43 responsible for the reduction in the precipitate free zone (PFZ). 44 In general, Sn has been considered as an effective element 45 which can interact with vacancies present in Al matrix and 46 also expected to change the precipitate morphology near the 47 grain boundaries by influencing the precipitate density and 48 interfacial energy between the precipitate and the matrix. 49 Presence of small and highly dense dispersoids improvised 50 by the minor addition of alloying elements can improve the 51 strength of Al alloys [10]. Stability in grain structure is com52 posed by delaying the recrystallization and pinning the grain 53 boundaries by the dispersoids, so appeared. Therefore, there 54 has been a strong motivation in adding Sn to 7075 alloy in 55 order to alter the dispersoids distribution pattern and as a con56 sequence retard the recrystallization process. Additionally, 57 the secondary phases (like ␩, T (Al2 Mg3 Zn3 ), S, ␪, (Al7 Cu2 Fe),) 58 formed during casting, may subsequently develop microseg59 regation [11–13]. The addition of alloying elements may also 60 cause the generation of more eutectic phases in the alloys. 61 Presence of high amount of coarse residual phases is proved 62 to be detrimental for toughness, fatigue strength, and stress 63 corrosion cracking (SCC) resistance properties for the alloy 64 [14–16]. 65 Restricting formation of dendritic structure and promoting 66 dissolution of the coarse residual constituents are accompa67 nied by the homogenisation treatment [1,13]. Additionally, 68 this process improves the ductility of the as-cast alloy by 69 enhancement of the plastic domain before failure. Develop70 ment of crystallographic texture or in other words orientation 71 of the grains becomes inevitable due to possible recrystal72 lization or grain growth phenomenon associated with the 73 homogenisation treatment [17–19]. Furthermore, phase trans74 formation and formation of different eutectic phases also alter 75 crystallographic texture [20]. Therefore, bulk texture anal76 ysis by laboratory X-ray machine equipped with a texture 77 goniometer becomes essential for in-depth understanding of 78 the microstructure [21–23]. 79 A high density of fine dispersoids can be formed dur80 ing homogenisation treatment which has a great effect on 81 grain size, recrystallization, and mechanical properties. The 82 main effect of dispersoids is to obstruct the recrystallization 83 behaviour during thermo-mechanical treatment. The effec84 tiveness of dispersoids can be calculated by ZPP (Eq. 1) which 85 is a good indicator for measuring the influence of a dispersoid 86Q4 on materials properties: 30 31

87

Z=k

 fy  r

(1)

Table 1 – Chemical compositions of the investigated alloy (in wt.%). Zn

Mg

Cu

Cr

Fe

Si

Mn

Sn

5.18

2.37

1.62

0.12

0.18

0.14

0.023

0.32

Al Bal.

Where k is a scalling factor, f is a volume fraction of dispersoids particles, y is the boundary energy, r is the average radius of the dispersoids. The value of ZPP can be tailored either by volume fractions and radius of dispersoids. Recrystallisation will be accompanied for regions with local fluctuations of ( f/r) ratio that falls below the critical value and vice versa [24]. The ZPP value will be enough at a critical value of ( f/r) ratio to stabilize the microstructure by overcoming the driving force for grain boundary migration and delay in the recrystallization process. Ling [10] explored that, the regions in microstructure where the values for local ( f/r) ratio falls below the critical value, will recrystallized. By contrast, the regions where the local ( f/r) ratio go beyond the critical value will not recrystallized. The size and distribution of dispersoids are very much affected by homogenisation time and temperature. The reasons stated above laid the foundation for examining the evolution of eutectic phases, dispersoids distribution, and homogenisation mechanism. It also aims at investigating the microstructural and crystallographic texture development for Al-Zn-Mg-Cu alloy containing a minor amount of Sn during the cast and homogenised conditions. The distribution of the dispersoids and the change of ( f/r) value with homogenisation time has been evaluated by TEM. The formation of different eutectic phases of the cast and homogenised alloys were also studied by different characterizing tools.

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Materials and methods

The investigated alloy prepared by the chill casting method bears the chemical composition given in Table 1. The temperature for homogenisation was selected according to the results obtained from differential scanning calorimetry (DSC) analysis. Cast samples were homogenised at 465 ◦ C for 0–48 h followed by quenching in air. The heating rate and the accuracy of the temperature for the homogenisation process were maintained by a PID controller (±1 ◦ C). DSC (Pyris Diamond, USA) was used with a constant heating rate of 10 ◦ C min−1 to identify the melting temperatures of the intermetallic phases. Specimens used for DSC investigation were mechanically thinned to reduce the weight to 11 mg. High purity Cu pans were utilized as the reference pans in this work. Microstructure and phase observation were jointly performed by optical microscopy (OM) and scanning electron microscopy (SEM). The chemical composition of intermetallic phases was revealed by SEM based energy dispersive X-ray Spectroscopy (EDS). Specimens for OM and SEM were polished following the traditional metallographic process and subsequently etched by Keller’s reagent [9]. In SEM, the secondary electron (SE) mode has been used with JEOL JSM-7600F FESEM and HITACHI S-3400N, operated at 15 kV. The morphology of the dispersoids and their frequency distribution were measured by transmission electron microscopy (TEM). For TEM investigations, samples were mechanically

Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055

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thinned down to 70 ␮m and then 3 mm diameter disc was extracted from the thin foils. The TEM samples were electropolished by using 30% HNO3 in methanol solution kept at −25 ◦ C, operated at a voltage of 35 kV. TEM samples were examined by JEM2200FS/Cs, operated at 200 kV. The identification of the eutectic phases of the cast and homogenized samples has been done by X-ray diffraction (XRD-Bruker D8 advance), operated at 35 kV/25 mA combinations and at a scan rate of 0.02◦ /s under Cu K␣ radiations. A texture goniometer (Bruker D8 Discover) with CuK␣ radiation was incorporated for bulk texture measurement. The details of the texture measurement and description of the sample orientations have been mentioned elsewhere [25].

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Results and discussion

3.1.

Microstructure after casting

Exhibition of (Fig.1a, b) serious dendritic segregation due to non-equilibrium solidification has been the most common and widely reported phenomenon occurs during the casting operation. The average grain size of the cast samples was calculated from SEM micrographs and recorded as 45 ␮m (±3 ␮m). Lamellar shaped low melting eutectic phases (Fig. 1b) were also revealed in this material. Between the two kinds of eutectic phases noticed in the cast sample, the chemical composition of the white portion of the eutectic phase exhibited a nominal composition (in at %) Al-43.86, Mg-18.54, Zn-15.74, Cu-12.23, which is close to (␩) Mg (Zn, Cu, Al)2 phase (Fig. 1c) [26,27], revealed by SEM-EDS. Interestingly, the phase Mg (Zn, Cu, Al)2 has the same crystallographic prototype as (MgZn)2 [23,28]. The grey portion of the eutectic structure of Al2 Cu (␪ phase) [23,25] bearing composition (at %) Al-59.82, Cu-33.81, Mg-2.75, Zn-2.34, Sn-1.28 was shown in (Fig. 1d). Therefore, it can be concluded that as-cast alloy microstructure consisted of a primary solid solution of ␣ (Al) and mainly two kinds of eutectic phases (a) white portion consisting of (␩) Mg (Zn, Cu, Al)2 phase and (b) copper-rich grey portion made of Al2 Cu (␪ phase). However, available literature revealed the presence of some other kinds of intermetallic phases like- Al7 Cu2 Fe, T (Al2 Mg3 Zn3 ), Mg2 Si, Mg32 (Al Zn)49 depending on the chemical compositions, solidification rates, and variation of casting process [1,11,13]. From the composition elemental maps of the Cu, Sn, Zn to Mg (Fig. 2) in as-cast material it can be observed that the concentrations of the elements were mainly coagulated close to the grain boundary with a gradual decrease towards the grain interior. During the homogenisation process, the atoms are diffused across the concentration gradient under the influence of thermal energy. The relationship between diffusion coefficient and temperature can be written as [17], Q D=D0 exp-( ) RT

(2)

Where D0 , Q, R, and T are diffusion co-efficient, diffusion activation energy, gas constant, and temperature in absolute scale respectively. It is well known that the homogenisation temperature has the most intense effect on the diffusion rate and its co-efficient. The necessity for optimizing the homogenisation

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temperature lies in the fact that too high a temperature may initiate burning of the cast sample.

3.2.

Phase transformations during homogenisation

The secondary (SE) images and corresponding EDS for 6 h homogenised alloys were represented in Fig. 3(a–e). Two distinct eutectic phases were formed after 6 h of homogenisation. The first of its kind was enriched with Cu and Mg, bearing composition (at%) of Al-46.61, Mg-18.46, Cu-22.46, Zn-2.39, Sn-1.08 (Fig. 3b, c) and its stoichiometry was close to Al2 CuMg (S phase) (marked in arrow). The second eutectic phase containing a high amount of Sn and Cu, bearing composition (at%) of Cu63.61, Sn- 22.46, Mg-3.73 Al- 2.38, Zn-1.08 (Fig. 3d,e) and close to Cu3 Sn (marked in arrow). Interestingly, the S and Cu3 Sn phases were absent in the cast material. A phase transformation from ␩ to S phase was reported in the homogenised (6 h) sample [13,25]. It can be seen from Fig. 2 that a high amount of different solute segregation has occurred in the inter-dendritic eutectic zone populated with Cu, Mg, Zn, Sn elements during casting of the alloy. The solute atoms exhibited preferential diffusion from the eutectic phase towards the Al matrix. Among all the solute atoms, the diffusion rate of Cu is the lowest compared to that of Zn, Mg, or Sn which understandably developed a concentrated pool of Cu in this region. Cu being the slowest diffusion species and hence rate-controlling, it became responsible for nucleation of the S phase [13,29]. Therefore, the driving force for the phase transformation from ␩ to S is supersaturation of Cu and subsequent development of stable S phase [30]. The amount of Zn in the S phase was less compared to primary ␩ phase indicating the fact that Zn atoms diffused faster from ␩ phase after 6 h of homogenisation Fig. 3(c) [13,28]. It can also be notated that S phase grew along with the ␩ phase and the growth of S phase is dependent mainly on the diffusivity of Cu, Mg atoms and inter-diffusion between the ␩ phase and the Al-matrix. On the other hand, a high concentration of Sn along the grain boundary was also held responsible for the formation of Cu3 Sn phase (Fig. 3d, e). After 6 h of homogenisation, both S and Cu3 Sn phases gradually grew in size and after 12 h of homogenisation, both phases finally converted to round or elliptical morphology Fig. 4(a). This was due to ripening and subsequent coalescence of S and Cu3 Sn phase. With a further length of homogenisation (24 h), the eutectic phases were gradually dissolved into the matrix [11] Fig. 4(b) and finally heading towards a stage (48 h) where all the intermetallic phases were almost dissolved into the matrix making the coarse particles extremely fine and dispersed (Fig. 4c, d) all over. The particles formed at this stage bears the chemical composition (at %) Al-87.54, Mg-1.48, Zn2.48, Cu-2.69, Sn-5.81 (Fig. 4e). It can be seen that fractions of residual phases were gradually reduced with increasing the homogenisation time (Fig. 4f).

3.3. DSC

Intermetallic phase analysis through XRD and

Fig.5(a,b) illustrated the differential scanning calorimetry (DSC) curves for as-cast and homogenized alloys. The two endothermic peaks appeared at 470◦ and 612 ◦ C in cast sample indicated the melting points of the eutectic phases and that

Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055

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Fig. 1 – SEM micrographs of the as-cast microstructure (a) low magnification (b) high magnification and corresponding their EDS results of the eutectic phases (c) Mg (Zn,Cu,Al)2 (d) Al2Cu.

Fig. 2 – SEM microstructure and main chemical elemental distribution maps in cast condition (a) back scattered image (b) Al (c) Cu (d) Zn (e) Mg (f) Sn.

Fig. 3 – SEM images of homogenised alloys at 465 ◦ C for (a) 6 h (b,c) high magnification view of the area of S phase and its EDS spectrum and (d,e) Cu3Sn phase and its EDS spectrum.

Fig. 4 – SEM micrographs of homogenised alloys at 465 ◦ C for (a) 12 h (b) 24 h (c) 48 h (d,e) high magnification view of residual phase and EDS spectrum (f) variation in the area fraction of eutectic phase with homogenisation time.

Fig. 5 – DSC curves of (a) as the cast and (b) homogenised alloy (c) XRD patterns of the alloy under as cast and different homogenized conditions.

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of the alloy respectively (Fig. 5a). This result suggested that the homogenisation temperature should not exceed 470 ◦ C to avoid possible danger of crossing non equilibrium solidus tem-

perature [25]. The disappearance of the peak at 470 ◦ C after homogenisation for 6 h suggested the dissolution of the nonequilibrium ␩ phase into the matrix (Fig. 5b). The appearance

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of a new endothermic peak at 488 ◦ C may be attributed due to the presence of copper-rich non equilibrium intermetallic phase S (Al2 CuMg) after homogenisation for 6 h [13,25]. The results were further corroborated by XRD and SEM analysis. After 48 h of homogenization treatment, the absence of any endothermic peak bore the signature of complete dissolution of the eutectic phases. Final homogenisation for 48 h almost completely dissolved the non-equilibrium eutectic phases into the matrix, generated during casting and evident from DSC plot (Fig. 5b). X-ray diffraction patterns of the as-cast and homogenised alloys (0–48 h) were presented in Fig. 5(c). The intermetallic phases in the as-cast sample mainly consisted of ␣ (Al), ␩ (MgZn2 ), Al7 Cu2 Fe, ␪ (Al2 Cu) phases. The small amount of Al7 Cu2 Fe phase was not detected from SEM microstructures due to its non-uniform presence throughout the Al matrix. After 6 h of homogenization treatment, a phase transformation from ␩ (MgZn2 ) to mainly Cu and Mg-enriched non-equilibrium aluminides S (Al2 CuMg) was noticed. These results were very much consistent with the earlier reports [14,15,31]. It was important to note that the intensities of the other peaks for ␩, ␪ phases were reduced and the phase ␪ completely disappeared after 6 h heat treatment. The eutectic phases of S (Al2 CuMg) and (Cu3 Sn) were still present even after 12 h of homogenization treatment. The dissolution of intermetallic phases gained momentum after 24 h of homogenisation, and the process reached nearing completion at 48 h, leaving only the peaks of Al7 Cu2 Fe and ␣ (Al). The phase Al7 Cu2 Fe remained unchanged during the whole homogenization treatment which is indicative of small dissolvability of Fe rich particles in the Al matrix. Therefore, the homogenization treatment did not leave a significant effect on the size and morphology of Fe-bearing secondary phases, a similar report has been presented in the work of Fan et al. [13].

3.4. Effect of homogenisation on dispersoids distribution Presence of uniformly distributed dispersoid particles during homogenisation is well known for their role in delaying recrystallization [10,23]. Undoubtedly, the shape, size, and distributions of the dispersoids are greatly affected by homogenisation temperature and time [17]. TEM images of dispersoids particles (MgZn2 ) within the grains after different time intervals of homogenisation (6–48 h) [32] were shown in Fig. 6(a–c). It was clearly indicated that with increasing the homogenisation time, the average size of the dispersoids was increased. During homogenisation, solute-rich atoms diffused to solute lean regions, increasing the diffusion rate of the solute atoms with time. After 6 h homogenization, fine dispersoids with high density were formed as shown in Fig. 6(a). Further, prolong homogenization would decrease the volume fractions of the dispersoids and increased the average particle size following Ostwald ripening mechanism (Fig. 6b,c). Fig. 6(d) exhibited dispersoids sizes and ( f/r) ratio for different homogenisation times. ZPP has remained an important parameter or indicator of ( f/r) ratio, which actually controls the pinning force on the grain and the subgrain boundaries

[10]. Higher homogenisation time accelerated the coarsening of the dispersoids and simultaneously lowered ZPP.

3.5.

Estimation of homogenisation kinetics

(3)

¯ average concentration of the elements Where, C: L: interdendritic spacing A0 : initial amplitude of the composition segregation and A0 =

1 1 (Cmax − Cmin ) = C0 2 2

(4)

The boundary conditions [ A (t) ] imposed by Fick’s second law is as follows [31]:

 A (t) = A0 exp

−

42 Dt L2



 −

(5)

42 D0 t exp L2

 Q   −

RT

A (t) 1 = A0 100

(6)

(7)

Then, −

42 D0 t exp L2

 

 Q −

RT

=

1 100

(8)

Further simplification of Eq. (8) leads to, R 1 = ln T Q



42 D0 t 4.6 L2

Assuming, A = 1 = A ln T

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321

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331



A0 exp

312

325

Eq. (5) can be simplified with the assumption that the elements are homogeneously distributed with their segregation level as low as 1% [34], i.e.,



311



Combining Eq. (1) and Eq. (4): A (t) = A0 exp

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Periodic arrangements of the main chemical elements along the interdendritic regions [25] invited the necessity for the evolution of diffusion kinetics, with respects to elements, during homogenisation. Fourier series components as a cosine function are popularly used to express the initial concentration of alloying elements: [31,33] 2x ¯ C(x) = C+A 0 cos L

308

 t  BL2

332

333

 (9)

R Q,

B=



4.6 42 D0

334

 ,

335

(10)

Eq. (10) is universally known as the homogenisation kinetics equation. The homogenisation curves can be obtained with the help of other parameters during casting. As the diffusion coefficient of Cu is the lowest among all other elements under investigation, the diffusion co efficient of Cu becomes rate-controlling [26,34]. The homogenisation kinetic curves

Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055

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Fig. 6 – TEM microstructures of the alloys homogenised for different time intervals (a) 6 h (b) 24 h (c) 48 h (d) size evolution of dispersoids and (f/r) ratio during different time for homogenisation.

Fig. 8 – ␸2 = 0◦ ,45◦ ,65◦ sections of ODFs of (a) as cast (b) homogenised sample for 48 h (c) volume fraction of different texture components.

Fig. 7 – Curves showing homogenisation kinetics.

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for Al-Zn-Cu-Mg-Sn sample can be achieved (Fig. 7) using D0 (Cu) = 0.084 cm2 /s, Q (Cu) = 136.8 KJ/mol and R = 8.314 J/mol. K in Eq. (10). It revealed that the holding time was decreased significantly with increasing the homogenization temperature, especially when the temperature goes beyond 450 ◦ C. Therefore, there were two paths to reduce the homogenisation time i.e. (a) increasing the homogenisation temperature or (b) grain size refinement. The most effective way to minimize the homogenisation time is to reduce the grain size by the addition of Sn in 7075 alloys, bearing the threat of overburning (as predicted by DSC results for cast sample) in mind. In the present study, the average dendritic spacing (L) was measured using quantitative metallography of the SEM micrographs and the value was found to be 50 ␮m. The kinetic model generated the optimum values for temperature (465 ◦ C) and soaking time (46.5 h), which commensurate with the experimental results.

However, the intensity of ␥ fibre was high and remained almost unchanged for both samples. The volume fractions of different texture components for both samples were shown in the histogram (Fig. 8c). It was clear that, Goss {011} <100 > , Brass {011} <211 > , P {011} <111 > , S {123} <634 > , CuT {552} <115> and Copper {112} <111> were the main components in the cast sample. On the contrary, homgenised materials contained strong Cube {100} <001 > , rotated Cube {110} <011> and moderate E, and F texture components [35]. The combined effect of high temperature and stress fields were held responsible for the formation of Brass, S and CuT texture components in the cast sample [25]. Formation of P-type of orientation during casting carried the footprint of the interaction between slip dislocation and particles which subsequently create strain field throughout the matrix [3,36]. However, the matrix strain has been reduced after 48 h of homogenisation, indicating that the dislocation strain field and particles were weak [37,38]. The origin of Goss texture suggested the presence of the high energy grain boundaries in the cast samples [23,25]. The development of strong Cube component in the homogenised material was formed as recovery progresses and larger grains are generated [36]. From the whole texture analysis, it was indicated that strong ␥ were intensified for both cast and homogenised conditions suggesting the higher ductility of the material [18]. The addition of a small amount of Sn in the alloy significantly improved the formability of the alloy even in cast condition, which was well consistent with the earlier works [9,39,40].

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3.6. Development of texture during casting and homogenisation of the alloy The crystallographic textures of the cast and homogenised specimens were studied by laboratory-based X-ray fitted with a texture goniometer. It was obvious form orientation distribution functions (ODFs) (Fig. 8a,b) that the cast samples possessed relatively strong ␣ (spread from Goss to rotated Goss) fibre compared to that for the homogenised sample.

4.

Conclusions

1 Exhibition of serious dendritic segregation and formation of secondary eutectic phases in the form of Mg (Zn, Cu, Al)2 , Al2 Cu and coarse Al7 Cu2 Fe during casting has been formed and gradually dissolved with the course of homogenisation.

Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055

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2 The homogenisation kinetic model corroborated strongly with the experimental results in terms of the distribution of the intermetallic phases homogenized at 465 ◦ C for 48 h. 3 Homogenisation for 6 h revealed the presence of two eutectic phases namely S (Al2 CuMg) and Cu3 Sn. The transformation from ␩ to S was mainly controlled by Cu due to it is higher concentration and lower diffusion rate compared to the other diffusing species. 4 TEM micrographs revealed that fine dispersoids with higher volume fraction were formed after 6 h of homogenisation. By increasing the homogenisation time, small size dispersoids dissolved into the matrix and the average size of other dispersoids were increased. Higher homogenisation time accelerated to the coarsening of dispersoids with increasing the average dispersoids sizes and lowering the ZPP. 5 Crystallographic texture analysis demonstrated the presence of strong ␥ fibre in both samples. Additionally, grains were affected by local deformation due to high temperature and stress fields that prevailed during the casting process and lead to the formation of strong Goss, Brass, P, CuT and Cu components.

Conflicts of interest 420 Q5

The authors declare no conflicts of interest.

Acknowledgment 421 Q6 422 423 424 425 426

The financial assistance received from the Govt. of India (UGCF.No-41-992/2012(SR)) to carry out this work is appreciatively acknowledged. The authors are thankful to Dr. Rajib Kalsar (IISc Bangalore) for the texture measurements of the investigated materials. The authors are also thankful to Mr. Badal Das for the casting of the samples.

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Please cite this article in press as: Ghosh A, et al. Phase transformation and dispersoid evolution for Al-Zn-Mg-Cu alloy containing Sn during JMRTEC 900 1–8 homogenisation. J Mater Res Technol. 2019. https://doi.org/10.1016/j.jmrt.2019.08.055