Evolution of Microstructure and Texture during Hot Deformation of a Commercially Processed Supral100

Transcription

1 J. Mater. Sci. Technol., 2012, 28(6), Evolution of Microstructure and Texture during Hot Deformation of a Commercially Processed Supral100 Y. Huang BCAST, Brunel University, Kingston Lane, Uxbridge UB8 3PH, UK [Manuscript received September 22, 2011, in revised form February 20, 2012] The microstructure and texture in a commercially processed Al-6 wt% Cu-0.4 wt% Zr (Supral100) aluminium alloy have been investigated after annealing and hot tensile straining at 450 C, using a field emission gun scanning electron microscope (SEM) and electron backscatter diffraction (EBSD). The microstructure of commercially processed alloy had a relatively large fraction of high angle grain boundaries (HAGBs) which were aligned parallel to the rolling direction, and a strong texture. Annealing at 450 C led to an increase in the fraction of HAGBs and to an increase in HAGB spacing and these changes were progressively enhanced by subsequent tensile deformation. The increasing fraction of HAGBs was due to the annihilation of low angle grain boundaries (LAGBs). A sharpening of texture during annealing was attributed to preferential textural growth, and the reduction of texture at higher tensile strains led to the development of superplastic behaviour. The present work supports the view that the evolution of the fine grain microstructure during the high temperature straining of Supral100 is primarily due to the accumulation of a large area of grain boundary during the initial thermomechanical processing, and does not involve any unusual restoration processes. KEY WORDS: Aluminium; Superplastic deformation; Microstructure; Texture; Electron backscatter diffraction 1. Introduction For polycrystalline solids to exhibit superplastic flow they must have equiaxed fine grain (<10 µm) microstructures, comprised mainly of high angle boundaries, which are stable at relatively high testing temperatures (>0.5 T m, where T m is the melting point in degrees Kelvin). For commercial Al alloys processed by the ingot route, the required microstructures are produced after thermomechanical processing (TMP) either by static recrystallisation prior to superplastic forming (SPF), or microstructural evolution during the early stages of SPF [1]. The former procedure is used for alloys such as AA5083 (Al-Mg- Mn) and AA7475 (Al-Zn-Mg-Cr), whereas AA2004 (Supral100) is made superplastic normally by the latter route. Recently, efforts have been paid to enhance the superplastic performance of the material Ph.D.; Tel.: ; Fax: ; address: yan.huang@brunel.ac.uk. by severe plastic deformation, which allows submicron grain structures to be developed in the material [2], although this route is expensive and applicable to limited small scales. Supral100 has a nominal composition in wt% of Al-6Cu-0.4Zr, and consists of an Al-Cu solid solution which contains a dispersion of very fine (<10 nm) ZrAl 3 precipitates and some CuAl 2 particles. During production the material is cast from a high superheat ( 780 C), and rapidly cooled to avoid the formation of coarse ZrAl 3 particles and obtain a high level of Zr in solid solution. The material is aged at 360 C to precipitate fine ZrAl 3 particles, solution treated at 50 C to take most of the copper into solid solution, and hot rolled to break down the as-cast structure. A subsequent warm/cold rolling schedule produces a heavily cold worked structure from which a fine microstructure evolves during the early stages of SPF at 450 C. After TMP the material is stabilised against static

2 532 Y. Huang: J. Mater. Sci. Technol., 2012, 28(6), recrystallisation by ZrAl 3 dispersoids, and the evolution of a fine grain superplastic microstructure with strain at 450 C is attributed to continuous dynamic recrystallisation. However, this phenomenological description does not imply the operation of any specific mechanism. It is usually assumed that after processing the microstructure consists of elongated, highly dislocated, pre-existing grains from which subgrains of low misorientation develop during the early stages of elevated temperature deformation or on annealing. Humphreys and Hatherly [3] have described mechanisms by which recrystallisation may occur dynamically. These include subgrain growth, subgrain coalescence, subgrain rotation or assimilation (accumulation) of dislocations into subgrain boundaries [4 8]. Ridley et al. [9] studied Al-Cu-Zr alloys and proposed that the structural evolution observed was consistent with strain induced geometrical dynamic recrystallisation. McNelley et al. [10,11] characterised the microstructure of Supral100 and reported that after TMP the material contained deformation bands of alternating lattice orientation, parallel to the rolling direction, which corresponded to the symmetric variants of the brass texture component {011}<211>. During annealing at 450 C, high angle grain boundaries (HAGBs) develop from the transition regions between the bands, and it was proposed that these contribute to the development of small equiaxed grains during subsequent hot/superplastic deformation. In favour of the TMP plus hot deformation route for exploiting the superplasticity of the material, which is cheaper, more efficient and commercially viable than other alternative routes such as severe deformation processing, the present work describes a study which examines microstructural evolution during hot deformation of commercially processed Supral100 sheet, in order to clarify the mechanisms leading to the establishment of superplastic behaviour in the material. 2. Experimental The alloy investigated in the experiment was Supral100 of nominal composition Al-6 wt%cu- 0.4 wt%zr supplied in the form of sheet by Superform, Worcester. The initial sheet thickness was 1.6 mm and the cast and heat treated alloy had been processed by a route which involved warm/cold rolling. Prior to tensile straining, the effect of annealing at various temperatures up to 450 C, for 1 h, on microstructure and texture was studied. Tensile specimens with a gauge length of 22 mm were machined from the sheet material with the tensile axis parallel to the rolling direction, and annealed at 450 C for 1 h to remove any mechanical damage. The specimens were deformed at 450 C at a constant strain rate of 10 3 s 1 to predetermined strains using an Instron tensile machine fitted with a high temperature testing facility. After deformation the specimens were cooled rapidly to ambient temperature. They were then sectioned longitudinally parallel to the RD-ND plane through the centre of the specimen and, after mechanical polishing and electropolishing, were examined metallographically. Additional measurements were also made on sections parallel to the RD-TD orientation. Detailed characterisation of the samples was made in a Camscan Maxim 2040 FEGSEM, by both backscattered and secondary electron imaging and electron backscatter diffraction (EBSD). EBSD maps were obtained using an HKL Channel acquisition system and processed using Vmap, an in-house software development. The positions of HAGBs (>15 ) and low angle grain boundaries LAGBs) (1 15 ) in the EBSD maps were determined, the lower angular limit being selected to remove noise from the data as discussed by Humphreys [12]. The maps presented in this paper show the HAGBs as black lines. The LAGBs are not shown for clarity. The amounts of various identified texture components were determined from the EBSD maps, and a map pixel was deemed to have the specified texture if its orientation was within 15 of the ideal. 3. Results and Discussion 3.1 As-received material Fig. 1 is an EBSD map of the as-received microstructure of the as-received bulk material in the RD-ND plane, for the 1.6 mm thick sheet. Regions containing larger CuAl 2 particles are filled black, and for clarity, only the HAGBs are marked. A highly elongated or banded structure aligned parallel to the rolling direction is seen and is consistent with the heavy deformation that the material has received. It may be noted that the distribution of the large second phase particles and of the grain boundaries shows some heterogeneity. EBSD measurements showed that the heavily worked structure contained a high proportion of HAGBs and that the material had a texture comprised mainly of brass {011}<211> and S {123}<634> components, as seen in Table 1. These Fig. 1 EBSD map of the as-received microstructure from the RD-ND plane. Only high angle boundaries (>15 ) are shown for clarity

3 Y. Huang: J. Mater. Sci. Technol., 2012, 28(6), Table 1 Parameters of microstructure and texture from EBSD measurements Status of Plane HAGB spacing/µm Aspect HAGB Texture processing ND TD RD ratio /% Brass S As received RD-ND As received RD-TD Annealed RD-ND (350 C, 1 h) Annealed RD-ND (450 C, 1 h) RD-TD % RD-ND % RD-ND % RD-TD % RD-ND Fig. 2 EBSD maps for ND-RD plane sections of material annealed at 450 C for 1 h after tension for different strains. Only high angle boundaries (>15 ) are shown for clarity. (a) 0%; (b) 15%; (c) 65%; (d) 100% are typical rolling texture components for rolled aluminium, the brass component usually being particularly strong after high temperature deformation [13]. It is also seen in Table 1 that the HAGB spacing in the RD direction, measured for the RD-TD and RD-ND orientations, shows very good agreement, as does the %HAGBs. 3.2 Static annealing of the as-processed sheet material Observations on the RD-ND plane (through thickness) showed that annealing for 1 h at 350 C and 450 C led to recovery involving an increase in the fraction of HAGBs, to an increase in HAGB spacing in both the RD and ND directions and to a reduction in aspect ratio (RD/ND) (Table 1). No recrystallisation was observed. There was also a slight sharpening of the brass and S textures (Table 1). Fig. 2(a) is an EBSD map of material annealed for 1 h at 450 C which shows the high density of HAGBs and some heterogeneity of the microstructure. Measurements made on an RD-TD (rolling plane) section are included in Table 1 and show that after annealing at 450 C for 1 h the grains have a roughly equiaxed shape with an aspect ratio (RD/TD) of The dimensions in the RD direction for both orientations are relatively similar i.e. 4.4 µm and 5.0 µm. 3.3 Microstructural evolution during deformation at 450 C Fig. 2(b) (d) show the evolution of microstructure as a function of tensile strain; only HAGBs (>15 )

4 534 Y. Huang: J. Mater. Sci. Technol., 2012, 28(6), Fig. 3 Evolution of boundary misorientation as a function of elongation for hot deformation at 450 C: (a) 0%, (b) 15%, (c) 65% and (d) 100% are included in these figures. There is a progressive increase in grain dimensions along both ND and RD directions, accompanied by a fall in aspect ratio (Table 1). It is noted that there is still some heterogeneity of grain size at the highest strain, but EBSD analysis shows that this is not associated with any specific textural component and that, overall, the grain orientations are essentially random. The analysed results of EBSD measurements showing the change of misorientation distribution with elongation are given in Fig. 3. It is seen that the proportion of HAGBs increases with increasing deformation, and reaches almost 80% at 100% elongation (Table 1). 3.4 Mechanism of microstructural evolution The as-received heavily worked sheet material contained a relatively high proportion of closely spaced HAGBs and had a strong texture comprised mainly of brass and S components. Substantial recovery occurred after annealing for 1 h at 350 C and 450 C, resulting in an increase in the proportion of HAGBs and their spacing in both the RD and ND directions, and a slight sharpening of the brass and S components of texture. Since the increase in spacing is essentially the result of grain growth, the area of grain boundary per unit volume will decrease. Hence, the increase in the proportion of HAGBs observed must result from the annihilation of LAGBs by recovery processes during annealing. Subsequent tensile deformation at 450 C, of specimens which had been annealed at that temperature, resulted in a progressive increase in grain dimensions in the RD and ND directions, a fall in RD/ND aspect ratio and a continual lessening of texture. Measurements made on the rolling plane (RD-TD) showed that the grain dimension in the TD direction did not undergo much change with strain, while the RD values were in reasonable agreement with the corresponding values obtained for the RD-ND orientation. As with annealing, it is seen that hot deformation is associated with grain growth (Fig. 2), such that the area of HAGB per unit volume decreases with increasing tensile strain. Analysis of EBSD data from which Fig. 2 was constructed leads to the boundary misorientation histograms seen in Fig. 3. These show that the proportion of HAGBs increases as LAGB s are continually removed from the microstructure. At 100% elongation most of the boundaries are HAGBs (Table 1). Simple calculations show that sufficient HAGBs existed in the initially processed sheet material to account for the HAGBs content after annealing and tensile straining. Approximating grain shape to that of a rectangular box, with dimensions equal to the boundary spacing in Table 1, it can be shown that the area of HAGB per unit volume is as follows: (a) as-received processed sheet: µm 2 /cm 3 ; (b) annealed at 450 C for 1 h: µm 2 /cm 3 ; (c) 65% elongation: µm 2 /cm 3 ; (d) 100% elongation: µm 2 /cm 3. It should be noted that since the TD spacing showed little change between the annealed at 450 C and 65% elongation conditions (Table 1), a value of 5.1 µm has been adopted

5 Y. Huang: J. Mater. Sci. Technol., 2012, 28(6), Fig pole figures for elongations of (a) 0%, (b) 15%, (c) 65% and (d) 100%, showing texture evolution during tensile deformation at 450 C for RD-ND plane. The maximum contour levels are 6.6, 6, 2.8 and 2.1 times of random, respectively for the TD spacing at 100% elongation. The above observations support the view that all of the HAGBs required to account for microstructural evolution and formation of an equiaxed fine grained microstructure on annealing and after tensile deformation already pre-exist in the material, as a result of the initial thermomechanical processing. These observations are consistent with the mechanism of strain enhanced geometrical dynamic recrystallisation [2,14]. The postulation of mechanisms for the generation of further HAGBs would appear to be unnecessary. 3.5 Texture evolution The brass and S textures observed in the initial material progressively weakened with increasing strain, particularly so at the higher strains (Table 1). Evolution of texture is shown in Fig. 4. It can be seen that the initially strong {011}<211> brass texture of the starting material decreased significantly during hot deformation. Such a weakening of the texture during superplastic deformation is frequently found [15,16], although the reasons for this are disputed. It is often assumed that the primary deformation mechanism during superplastic deformation is grain boundary sliding, and that the consequent quasi-random rigid body rotation of grains is responsible for the observed texture weakening [15]. However, it has also been suggested that slip deformation in fine-grained metals with a high strain rate sensitivity may be the primary mechanism of superplastic deformation [17], and that this gives better agreement with the present experimental results. 4. Conclusions An investigation into the effect of annealing and hot tensile deformation on the evolution of microstructure and texture in a commercially processed Al-6 wt% Cu-0.4 wt% Zr (Supral100) sheet material, using high resolution EBSD studies, led to the following observations:

6 536 Y. Huang: J. Mater. Sci. Technol., 2012, 28(6), The starting material had a strongly banded microstructure aligned parallel to the rolling direction, contained a relatively high proportion of HAGBs, and had a strong texture comprised mainly of Brass and S components. Annealing at 350 C and 450 C led to recovery involving an increase in the fraction of HAGBs, an increase in HAGB spacing, a reduction in aspect ratio (RD/ND), and to a sharpening of texture. No recrystallisation was observed on annealing. The increase in the proportion of HAGBs on annealing was attributed to the annihilation of LAGB s by recovery processes and the sharpening of texture to preferential textural growth. During tensile deformation of annealed specimens at 450 C there was a progressive increase in grain dimensions in the RD and ND directions, a decrease in the RD/ND aspect ratio, a reduction in texture, and an increase in the fraction of HAGBs. At 100% tensile elongation, the proportion of HAGBs reaches 80% due to the progressive removal of LAGB s and the microstructure consists of equiaxed grains of 5 7 µm in diameter, whose orientations are essentially random. The randomisation of texture observed at the higher strains is typical of superplastic behaviour which is characteristic of randomly orientated fine grain microstructures containing a high proportion of HAGBs. Simple calculations show that the HAGBs are required to account for microstructural evolution on annealing, and after subsequent tensile straining, they pre-exist in the material as a result of the initial thermomechanical processing. Acknowledgements The author is grateful to EPSRC for financial support via Grant GR/R69952/01 and to Superform for the supply of material. REFERENCES [1 ] J. Pilling and N. Ridley: Superplasticity in Crystalline Solids, Institute of Metals, London, 1989, 17. [2 ] R.B. Figueriredo, M. Kawasaki, C. Xu and T.G. Langdon: Mater. Sci. Eng. A, 2008, 493, 104. [3 ] F.J. Humphreys and M. Hatherly: Recrystallization and Related Annealing Phenomena, 2nd edn, Elsevier, Oxford, 2004, 388. [4 ] N. Nes: Metal Sci., 1979, 13, 211. [5 ] E. Hornbogen: Metall. Trans. A, 1979, 10, 947. [6 ] R. Gudmundsson, D. Brooks and J.A. Wert: Acta Metall. Mater., 1991, 39, 19. [7 ] S.J. Hales, T.R. McNelley and H.J. McQueen: Metall. Trans. A, 1991, 22, [8 ] R. Fernández, M. Mabuchi, K. Higashi and G. González-Doncel: Comp. Sci. Technol., 2009, 69, 373. [9 ] N. Ridley, E.M. Cullen and F.J. Humphreys: Mater Sci. Technol., 2000, 16, 117. [10] M. Eddahbi, T.R. McNelley and O.A. Ruano: Metall. Mater. Trans. A, 2001, 32, [11] T.R. McNelley, D.L. Swisher and M.T. Perez-Prado: Metall. Mater. Trans. A, 2002, 33, 279. [12] F.J. Humphreys: J. Mater. Sci., 2001, 36, [13] P.S. Bate and Q. Oscarrson: Mater. Sci. Technol., 1990, 6, 520. [14] H.J. McQueen, O. Knustad, N. Ryum and J.K. Solburg: Scripta Metall., 1985, 19, 73. [15] K.A. Pabmanabhan and K. Lücke: Z. Metallkd., 1986, 77, 765. [16] P.S. Bate, N. Ridley and B. Zhang: Acta Mater., 2007, 55, [17] J.L. Song and P.S. Bate: Acta Mater., 1997, 45, 2747.