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strength of Al-Cu binary alloys
Materials Science & Engineering A 707 (2017) 58–64
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Sc and Zr additions on the microstructure/strength of Al-Cu binary alloys
MARK
Thomas Dorina,⁎, Mahendra Ramajayama, Justin Lambb, Timothy Langanc a b c
Deakin University, Institute for Frontier Materials, Geelong, Victoria 3216, Australia Universal Alloy Corporation, USA Clean TeQ Ltd., Australia
A R T I C L E I N F O
A B S T R A C T
Keywords: Scandium Precipitation Recrystallization Dispersoids Aluminium
Scandium increases the strength of aluminium alloys via three mechanism: 1) solid solution strengthening, 2) precipitation hardening, 3) grain structure control. Despite the well documented benefits, scandium has found very limited use in commercial grade aluminium alloys due to its high cost. However, new efficient extraction technologies promise an ensured supply of scandium and a significant drop in cost. Development of the next generation of aerospace Al alloys will come from innovation in alloy chemistry. One such innovation could be the addition of scandium in combination with Zirconium which increases the specific strength and stiffness of aluminium alloys through the precipitation of the L12 Al3(Sc, Zr) dispersoid. However, very little is known about the interactions of the Al3(Sc, Zr) dispersoid and the θ′-phase. Here, the effects of Sc and Zr additions to a model Al-Cu alloy were examined. The precipitates were investigated through TEM and APT. EBSD was used to characterise the texture of the studied alloys. Finally, the ageing response of the alloys was monitored through tensile testing. The refinement of the Cu precipitates accounted for an increase of up to 120 MPa of the peak aged strength and the core/shell dispersoids accounted for up to 40 MPa.
1. Introduction Because of their high strength to weight ratio, aluminium alloys are a prime choice for use in structural aerospace applications. One of the main strategies to strengthen aluminium alloys is to micro-alloy with elements that result in the formation of a high number density of nanosize precipitates [1]. These precipitates act as obstacles to dislocation motion and hence improve an alloy's strength. In order to obtain the optimal mechanical properties, the precipitates’ distribution has to be carefully controlled, in terms of size, volume fraction and inter-particle spacing. The 2xxx-series of aluminium alloys, commonly utilised in the aerospace industry, are alloyed with Cu which results in the formation of hardening precipitates and Mn and/or Zr as the main texture control elements [2]. With the emergence of composite materials, the aerospace industry is investigating the use of scandium (Sc), in combination with Zr, as Sc is known to provide additional strength through solid solution strengthening, precipitation hardening, and grain structure control [3,4]. The hardening potential of Sc was studied in the 1970s with a large part of the work conducted on binary Al-Sc [5,6] and ternary Al-Mg-Sc
⁎
systems [4,6–9]. Extensive reviews of the scientific literature on Sc in aluminium alloys were conducted in 1998 [3], 2005 [10] and more recently [11]. The strengthening possible from Sc in Al is very unique (and beneficial), and disproportional to the level of alloying when compared to other elements [12,13]. Despite the low solubility of Sc in Al, it achieves exceptional strength whilst retaining reasonable ductility. The addition of Zr modifies the Sc-dispersoids by both refining their size and increasing their thermal stability [14–16]. As such, the combined addition of Sc and Zr in aluminium results in alloys with exceptional strength and improved recrystallization resistance [3,14,17–21]. The Sc atoms were also reported to form ternary AlCuSc phases [22], if carefully designed heat treatment schedule to favour the formation of Al3Sc and θ′ was not utilised [23]. Despite these benefits, Sc has found almost no commercial applications so far because it was previously considered to be prohibitively expensive. However, with the projected price reduction of Sc in the next few years, there is a recent regain of interest for these alloys. The strength contributions from Sc and Zr additions are multi-fold: 1) solid solution strengthening, 2) precipitation of Al3(Sc,Zr) and 3) grain structure control. Sc atoms were recently reported to act as
Corresponding author. E-mail address: [email protected] (T. Dorin).
http://dx.doi.org/10.1016/j.msea.2017.09.032 Received 12 July 2017; Received in revised form 5 September 2017; Accepted 8 September 2017 Available online 09 September 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 707 (2017) 58–64
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heterogeneous nucleation sites for the θ′ phase and to then segregate at the θ′/matrix interface [24,25]. There is no available research on the effect of adding both Sc and Zr to an Al-Cu alloys. In this study, we examined the effect of alloying a binary Al-Cu alloy with different Sc and Zr contents. Four alloys were selected for the study as follows: Al-4 wt%Cu, Al-4 wt%Cu-0.2 wt%Sc, Al-4Cu-0.1 wt% Zr, Al-4 wt%Cu-0.1 wt%Sc-0.14 wt%Zr. The alloys were extruded, solution treated and aged. The effect of composition on precipitation and recrystallization was investigated using electron microscopy and atom probe tomography. The strength increment during ageing was monitored with tensile tests. It was found that the yield strength could be increased by up to 150 MPa with the suitable Zr and Sc content. An Al3Zr shell was found to form around the Al3Sc dispersoids. This paper is the first to experimentally observe the nucleation of θ′ precipitates on the core-shell dispersoids. The strengthening contributions are analysed and we find that the refinement of the θ′ precipitates accounts for the major part of the strength increase.
strain during the tests. At least 3 samples were tested under each condition to ensure reproducibility. 2.3. Microscopy The state of recrystallization was evaluated with electron backscattered diffraction (EBSD) measurements in a Jeol JSM 7800F scanning electron microscope (SEM) at 20 kV. The EBSD samples were sectioned and metallographically prepared using the usual methods. The final polishing step used was 5–10 min of polishing using colloidal silica. The precipitates distribution was characterised in a JEOL 2100F transmission electron microscope (TEM) equipped with energy dispersive spectroscopy (EDS). TEM samples were prepared by electropolishing in a solution of 33% nitric acid in methanol at −20 °C. Atom probe tomography (APT) was carried out on a LEAP 4000 h instrument at a pulse fraction of 200 kHz, a temperature of 60 K, and pulse fraction of 20%. Samples were prepared by electro-polishing in a micro-loop apparatus under an optical microscope in a solution of 5% nitric acid in ethanol at room temperature [27,28]. Electro-polished tips were given a final sharpen using Ga ions at 10 kV in a Quanta dual beam focused ion beam instrument.
2. Materials and experimental methods 2.1. Alloys preparation
3. Results
The alloys used for this study were cast using master alloys to ensure maximum homogeneity of the cast metals. Sc and Zr additions were made via standard master alloys compositions; Al-2 wt%Sc and Al-5 wt %Zr. It should be noted that Cu additions were made without the use of a master as it is known to dissolve easily into liquid Al [26]. Melting was conducted in a small induction furnace with a ~ 4 kg capacity for aluminium. The molten metal was kept at 710 °C for 15–20 min before being poured into a cylindrical steel mould. The cylindrical mould comprised a large steel base to provide directional solidification. The compositions of the alloys cast for this study can be seen in Table 1. The as-cast alloys were then homogenised to form Sc and Zr dispersoids and to dissolve the coarse intermetallic phases left over from the casting process. At least four homogenised billets were then extruded for each composition. The extrusion process was conducted on a specially designed module adapted to a traditional Instron apparatus. A small resistance furnace was used to pre-heat the billets prior to extrusion. The billet was then forced into a die at the selected temperature. A circular die of 5.3 mm diameter was used for all the extrusions conducted in this project. It should be noted no extrusion defects such as incipient melting or broken surface were observed in any of the resulting F-temper extrusions. The F-temper extrusions were then solution heat treated at 500 °C for 1 h before being stretched and aged at 160 °C for various times in order to form the hardening θ′ phase. The alloys were allowed to naturally age for 72 h prior to artificial ageing.
3.1. Strengthening kinetics In order to monitor the mechanical properties evolution during the ageing process at 160 °C, tensile tests were conducted on the four alloys at different artificial ageing times. Tensile specimens were aged for 1 h, 3 h, 10 h and 20 h at 160 °C. The tensile stress/ strain curves are displayed in Fig. 1(a)–(e) for respectively alloy 1 to alloy 4. A clear trend of simultaneous increase in strength and decrease in ductility with increasing alloying content can be observed. The yield stress, tensile stress and uniform elongation evolutions for all 4 alloys are reported in Fig. 2(a)–(c). The highest yield stress is reached after 10 h at 160 °C which is defined as the peak aged condition for all 4 alloys. The most significant increase in strength was observed in alloy 4 where the yield stress was 150 MPa higher than that of alloy 1. A drop in uniform elongation as a function of ageing time was reported for all 4 alloys. After 10 h at 160 °C, the uniform elongation ranged from 6% to 8%. 3.2. Recrystallization To evaluate microstructural differences, electron backscattered diffraction imaging was conducted on alloys 1, 2, 3 and 4 after the solution treatment of 1 h at 500 °C (see Fig. 3). Alloy 1 was observed to be fully recrystallised with a grain size of about 30 µm. Alloy 2 was also fully recrystallised but the grain size was refined with an average measured grain size of 20 µm. Alloys 3 and 4 displayed a strong extrusion texture with elongated grains. Whilst alloy 3 was partially recrystallised near the surface of the extrusion, alloy 4 was completely unrecrystallised. These observation are in agreement with previous studies that report a more efficient grain pinning from combined Sc and Zr additions as compared to the single elements.
2.2. Mechanical testing Tensile tests were carried out in an Instron 5967 equipment mounted with a 5 kN load cell. All tests were conducted at constant cross head velocity of 1 mm/min. For a 16 mm gauge length sample, this constant cross head velocity is equivalent to a strain rate of ~ 0.001/s. A non-contact video extensometer was used to measure the
3.3. Precipitation Table 1 Compositions of the four alloys under study.
Alloy Alloy Alloy Alloy
1 2 3 4
Al
Cu
Sc
Zr
Fe
Ti
Bal Bal Bal Bal
4 4 4 4
0 0 0.2 0.1
0 0.1 0 0.14
0.07 0.07 0.07 0.07
0.005 0.005 0.005 0.005
To examine the effects of Sc and Zr on the precipitation behaviour, two samples from alloy 1 and 4 in the peak aged condition were observed in transmission electron microscopy (TEM). The samples were observed in the <100> zone axis as two of the four θ′ variants can be observed edge-on in this orientation. It was found that the length and width of the θ′ plates were significantly decreased in the alloy that contained Sc and Zr (see Fig. 4). In 59
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Fig. 1. Tensile stress-strain curves for samples aged for 0 h, 1 h, 3 h, 10 h and 20 h at 160 °C for (a) Alloy 1 to (d) Alloy 4.
Fig. 6(c)).
alloy 1, the plates’ width was 197 nm on average and the thickness was 5–10 nm. In Alloy 4, the average plates’ width was 38 nm and thickness 2–5 nm. The θ′ plates size distributions was reported in the form of histograms, as shown in Fig. 4. The distributions both exhibit a lognormal shape. The presence of dispersoids was also observed, as shown in Fig. 5. The dispersoids were confirmed to contain both Sc and Zr and a coreshell type contrast was observed. In most instances, the dispersoids were observed to be in contact with a θ′ plate. This suggests that the dispersoids play a significant role in the nucleation and growth of θ′ which results in a refined distribution. In order to confirm the core-shell nature of the Sc and Zr containing dispersoids, alloy 4 was examined in atom probe tomography (APT), Fig. 6. Two dispersoids were present in the APT volume and the coreshell structure was clearly confirmed. The Sc rich core was observed to be about 5 nm diameter and the shell had a total diameter of about 10 nm. Fig. 6(b) reveals that the dispersoids are embedded in the centre of the θ′ precipitates which confirms that the dispersoids act as preferential nucleation sites for these precipitates. The composition evolution across the plates and dispersoids confirms that the dispersoids cores are Sc rich and that the shells contain both Sc and Zr (see
4. Discussion 4.1. Precipitation mechanisms Al-Cu alloys have been extensively studied and the main strengthening phase is known to be the metastable θ′ phase [29]. A traditional way to promote and refine these precipitates is by introducing additional nucleation sites which is traditionally performed by prestretching up to 6%. This strategy does result in an increase in strength but has the main disadvantage of decreasing the alloy's ductility and fatigue performance in the peak aged condition. Furthermore, in some processes, such as casting or forging, it is not possible to introduce cold deformation prior to ageing which decreases the potential hardening obtained from θ′ precipitation. Here, it was observed that dispersoids with a Sc-rich core and Zr-rich shell act as efficient nucleation sites for θ′ particles and could be used as an additional strategy to promote the formation of the main hardening phase with little or no cold deformation required. Additions of Sc and Zr are known to result in the formation of L12 dispersoids. These dispersoids traditionally form during solidification 60
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Fig. 2. Evolution of (a) yield stress (MPa), (b) ultimate tensile stress (MPa) and (c) uniform elongation (%) as a function of ageing time at 160 °C for the 4 alloys under investigation.
The θ′ precipitates observed in the peak aged temper for alloy 4 are clearly associated with the dispersoids. This suggests that the dispersoids act as preferential nucleation sites for these precipitates. The size of the θ′ precipitates was significantly different in the sample with
and high temperature homogenisation. In the present study, the alloys were homogenised prior to extrusion, hence, the dispersoids were already formed and stable. The combined additions of Sc and Zr were observed to result in precipitates with a Sc-rich core and Zr-rich shell.
Fig. 3. EBSD map at 100× magnification for respectively (a) Alloy 1 to (d) Alloy 4.
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Fig. 4. High magnification TEM micrographs showing Cu precipitates in (a) Alloy 1 and (b) Alloy 4. Note the difference in the scale bar. (c) and (d) Respective precipitates size histograms.
the θ′/matrix interface. We suggest that the Sc observed around the θ′ plates, in previous works [24,25], comes from the nucleation of θ′ on the Al3Sc dispersoids and then of a redistribution of the Sc atoms around the θ′ plates. In the present work, the presence of the Zr shell prevents Sc redistribution. Thus the Zr stabilizes the Al3Sc whilst keeping the refinement effect of the θ′ plates.
Sc and Zr as compared to the sample free from Sc and Zr. Indeed, the presence of Sc and Zr in alloy 4 refined the average diameter of the plates to 38 nm as compared to 197 nm with no Sc nor Zr in alloy 1. As a result, the dispersoids were found to help nucleate and stabilise small θ′ precipitates, preventing rapid coarsening of these precipitates. Contrarily to previous reports, there was no Sc segregation observed at
Fig. 5. TEM-EDS analysis of Alloy 4 (Al-Cu-Zr-Sc) aged at 160 °C for 10 h.
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Fig. 7. Strength contribution from the θ′ precipitates σ θ′ as a function of precipitate diameter as predicted from Eq. (2) and measured on alloys 1 and 4 from this study.
1
σ θ′ = 0.13MGα 2
(2)
where M = 3 is the Taylor factor, G = 28 GPa is the Al shear modulus, b = 0.286 nm is the magnitude of the Burgers vector and r0 = 0.572 nm is the inner dislocation cut-off radius. The evolution of σ θ′ as a function of precipitate diameter is plotted in Fig. 7. α = 7.4 was used as fitting parameter to get the best match with the empirical measurements. This value is smaller than what was reported elsewhere where α was found to be greater than 20 [30–32]. As a conclusion, the increase in strength contributions from the θ′ precipitates comes from the formation of much smaller precipitates in the presence of Sc and Zr. As such alloys 1 and 4 have average precipitates diameter of respectively 197 nm and 38 nm and this corresponds to a strength contribution from θ′ increasing from 50 MPa to 120 MPa.
Fig. 6. Atom probe tomography reconstruction of an alloy 4 sample aged at 160 °C for 10 h.
4.2. Strength contribution from θ′ The strength increment from the θ′ precipitates can be obtained by computing the strength increment during the final natural and artificial ageing. Indeed, for all the studied alloys, the θ′ precipitates form during the final ageing treatment, hence, the strength increment during ageing corresponds to the contribution to strength of the θ′ precipitates only. These contributions are summarised in Table 2. The alloys 1–4 have the exact same Cu content, hence we can hypothesize that the same volume fraction of θ′ precipitates should form. The major strengthening differences observed, hence only come from differences in precipitate size. The size of the θ′ precipitates can be described with two parameters, a diameter Dp and a thickness tp . Previous empirical and phase field modelling studies [30–32] have observed a proportionality relationship for these two parameters such that:
Dp = α × tp
0.079Dp 1 b [0.23 + 0.038α 2 Dp +0.007α ]ln ( ) r0 Dp
4.3. Additional strength contributions Apart from the refinement of θ′ precipitates, the most significant strengthening contributions appear to come from
(1)
1) precipitation hardening from the Al3(Sc,Zr) dispersoids 2) grain structure control
where α is a constant that depends on the alloy composition and heat treatment conditions. Zhu and Starke [33] developed a strengthening model for unshearable plate-like particles applicable to the θ′ precipitates. The variables in the initial model are the volume fraction fv , average diameter Dp and thickness tp of the precipitates. In the present study, the bulk composition in Cu is the same in the four studied alloys and if we assume that the stoichiometric volume fraction of fv = 0.051 is reached for all alloys after ageing to peak strength. As a result, Zhu and Starke's model can be computed as a function of the precipitate diameter Dp , and with one fitting parameter α , as follows:
To estimate the importance of these contributions, the tensile data of the as-quenched samples are used. As such, the strength of the asquenched alloy 1 (free from dispersoids) is used as a reference and is subtracted from the strength of the other samples as quenched. Indeed, all alloys contain the same amount of Cu so it can be considered that after the solution treatment, the only difference in the strength measured should come from the dispersoids and the resulting grain structure. The calculated strength increment are reported in Table 3. The observed strength increment is similar for alloys 2 and 3 with only Zr or only Sc which confirms that addition of Sc only does not provide
Table 2 Measured yield strength increment coming from the formation of θ′ precipitates.
Table 3 Measured yield strength increment coming from the L12 dispersoids and from grain size refinement for the 4 studied alloys.
Alloy
1
2
3
4
Alloy
1
2
3
4
σ θ′ (MPa)
50
74
76
120
σdisp (MPa)
/
17
20
40
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alloys used in this project. Deakin University’s Advanced Characterisation Facility is acknowledged for use of the Jeol JSM 7800F scanning electron microscope, Jeol 2100F transmission electron microscope and LEAP 4000 HR atom probe tomography instrument.
significant strength increase. However an extra 20 MPa is obtained by the addition of both Sc and Zr in alloy 4. The strength differences in the as-quenched condition is relative to the dispersoids’ volume fraction. Considering an Al3M stoichiometry and based on the alloying content from Table 1, the fraction of dispersoids should be similar in Alloy 3 and 4 which confirms that Al3(Sc, Zr) dispersoids are most potent hardener than the Al3Sc ones. Similarly, the stoichiometric fraction of dispersoids should be lower in Alloy 2 as compared to Alloy 3 and a similar strength contribution is observed in these alloys confirming again that the addition of Sc alone does not work efficiently in these alloys (Table 3). From this data, one realises that the most significant contribution to strength is by far the refinement of the θ′ precipitates (more dispersoids = more nucleation sites) rather than the direct hardening effects from the dispersoids or reduced grain size.
References [1] T. Gladman, Mater. Sci. Technol. 15 (1) (1999) 30–36. [2] E.A. Starke, J.T. Staley, Prog. Aerospace Sci. 32 (2-3) (1996) 131–172. [3] L.S. Toropova, Advanced aluminum alloys containing scandium: structure and properties, Taylor & Francis, New York, 1998. [4] K.L. Kendig, D.B. Miracle, Acta Mater. 50 (16) (2002) 4165–4175. [5] N. Blake, M.A. Hopkins, J. Mater. Sci. 20 (8) (1985) 2861–2867. [6] B.A. Parker, Z.F. Zhou, P. Nolle, J. Mater. Sci. 30 (2) (1995) 452–458. [7] R.R. Sawtell, C.L. Jensen, Metall. Mater. Trans. A 21 (1) (1990) 421–430. [8] Y.A. Filatov, V.I. Yelagin, V.V. Zakharov, Mater. Sci. Eng. A 280 (1) (2000) 97–101. [9] S. Lathabai, P.G. Lloyd, Acta Mater. 50 (17) (2002) 4275–4292. [10] J. Røyset, N. Ryum, Int. Mater. Rev. 50 (1) (2005) 19–44. [11] T. Dorin, M. Ramjayam, A. Vahid, T. Langan, “Chapter 17: aluminium scandium alloys,” in: Fundamentals of Aluminium Metallurgy: Recent Advances, Elsevier, 2017 (In press). [12] E.A. Marquis, D.N. Seidman, Acta Mater. 49 (11) (2001) 1909–1919. [13] E.A. Marquis, D.N. Seidman, D.C. Dunand, Acta Mater. 51 (16) (2003) 4751–4760. [14] V.I. Yelagin, V.V. Zakharov, S.G. Pavlenko, T.D. Rostova, Phys. Met. Metall. 60 (1) (1985) 88–92. [15] V.G. Davydov, T.D. Rostova, V.V. Zakharov, Y.A. Filatov, V.I. Yelagin, Mater. Sci. Eng. A 280 (1) (2000) 30–36. [16] C.B. Fuller, D.N. Seidman, D.C. Dunand, Acta Mater. 51 (16) (2003) 4803–4814. [17] V.G. Davydov, V.I. Elagin, V.V. Zakharov, T.D. Rostova, Metall. Term. Obra. Metallov. (1996) 25–30. [18] T.D. Rostova, V.G. Davydov, V.I. Yelagin, V.V. Zakharov, Mater. Sci. forum 331 (2000) 793–798. [19] K.E. Knipling, D.N. Seidman, D.C. Dunand, Acta Mater. 59 (3) (2011) 943–954. [20] J.D. Robson, Acta Mater. 52 (6) (2004) 1409–1421. [21] Y.W. Riddle, T.H. Sanders, Metall. Mater. Trans. A 35 (1) (2004) 341–350. [22] P. Riani, L. Arrighi, R. Marazza, D. Mazzone, G. Zanicchi, R. Ferro, J. Phase Equilibria Diffus. 25 (1) (2004) 22–52. [23] M. Nakayama, Y. Miura, Jpn Inst. Light Metals 46 (6) (1996) 275–279. [24] B.A. Chen, L. Pan, R.H. Wang, G. Liu, P.M. Cheng, L. Xiao, J. Sun, Mater. Sci. Eng. A 530 (2011) 607–617. [25] C. Yang, P. Zhang, D. Shao, R.H. Wang, L.F. Cao, J.Y. Zhang, G. Liu, B. Chen, J. Sun, Acta Mater. 119 (2016) 68–79. [26] V.S. Zolotorevsky, N.A. Belov, M.V. Glazoff, Casting Aluminum Alloys, 12 Elsevier, Amsterdam, 2007. [27] B. Gault, M.P. Moody, J.M. Cairney, S.P. Ringer, Atom Probe Microscopy, Springer, New York, 2012. [28] M.K. Miller, R.G. Forbes, Atom-Probe Tomography: The Local Electrode Atom Probe, Springer, New York, 2014. [29] J.M. Papazian, Metall. Trans. A 12 (2) (1981) 269–280. [30] A.W. Zhu, J. Chen, E.A. Starke Jr, , Acta Mater. 48 (9) (2000) 2239–2246. [31] A. Biswas, D.J. Siegel, C. Wolverton, D.N. Seidman, Acta Mater. 59 (15) (2011) 6187–6204. [32] H. Liu, B. Bellon, J. Llorca, Acta Mater. 132 (2017) 611–626. [33] A.W. Zhu, E.A. Starke, Acta Mater. 47 (11) (1999) 3263–3269.
5. Conclusions In the present study, the effects of adding Sc and Zr to an Al-4 wt %Cu alloy were investigated. Four alloys with different Sc/Zr ratio were analysed. The most interesting results can be summarised as follow:
• The yield strength was doubled for the alloy Al-4 wt%Cu-0.1 wt%Sc• • • • •
0.14 wt%Zr as compared to its binary counterpart. The peak strength was found to occur after 10 h of ageing at 160 ᴼC for all four alloys. Recrystallization was fully inhibited for the alloys containing both Sc and Zr. The Sc and Zr dispersoids were observed to form a core rich of Sc and a shell rich in Zr. The dispersoids helped the nucleation and stabilisation of θ′ precipitates during ageing. The θ′ precipitates were refined from an average diameter of 197 nm in alloy 1–38 nm in alloy 4. The contribution to yield strength from the θ′ precipitates was found to be up to 120 MPa. The contribution from the dispersoids + grain size was up to 40 MPa. Hence, the most important contribution to strength came from the role of the dispersoids in refining the θ′ precipitates.
Acknowledgements The authors would like to acknowledge Clean TeQ for providing inkind Al-Sc master alloys. Dave Gray is warmly thanked for casting the
64
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Sc and Zr additions on the microstructure/strength of Al-Cu binary alloys
MARK
Thomas Dorina,⁎, Mahendra Ramajayama, Justin Lambb, Timothy Langanc a b c
Deakin University, Institute for Frontier Materials, Geelong, Victoria 3216, Australia Universal Alloy Corporation, USA Clean TeQ Ltd., Australia
A R T I C L E I N F O
A B S T R A C T
Keywords: Scandium Precipitation Recrystallization Dispersoids Aluminium
Scandium increases the strength of aluminium alloys via three mechanism: 1) solid solution strengthening, 2) precipitation hardening, 3) grain structure control. Despite the well documented benefits, scandium has found very limited use in commercial grade aluminium alloys due to its high cost. However, new efficient extraction technologies promise an ensured supply of scandium and a significant drop in cost. Development of the next generation of aerospace Al alloys will come from innovation in alloy chemistry. One such innovation could be the addition of scandium in combination with Zirconium which increases the specific strength and stiffness of aluminium alloys through the precipitation of the L12 Al3(Sc, Zr) dispersoid. However, very little is known about the interactions of the Al3(Sc, Zr) dispersoid and the θ′-phase. Here, the effects of Sc and Zr additions to a model Al-Cu alloy were examined. The precipitates were investigated through TEM and APT. EBSD was used to characterise the texture of the studied alloys. Finally, the ageing response of the alloys was monitored through tensile testing. The refinement of the Cu precipitates accounted for an increase of up to 120 MPa of the peak aged strength and the core/shell dispersoids accounted for up to 40 MPa.
1. Introduction Because of their high strength to weight ratio, aluminium alloys are a prime choice for use in structural aerospace applications. One of the main strategies to strengthen aluminium alloys is to micro-alloy with elements that result in the formation of a high number density of nanosize precipitates [1]. These precipitates act as obstacles to dislocation motion and hence improve an alloy's strength. In order to obtain the optimal mechanical properties, the precipitates’ distribution has to be carefully controlled, in terms of size, volume fraction and inter-particle spacing. The 2xxx-series of aluminium alloys, commonly utilised in the aerospace industry, are alloyed with Cu which results in the formation of hardening precipitates and Mn and/or Zr as the main texture control elements [2]. With the emergence of composite materials, the aerospace industry is investigating the use of scandium (Sc), in combination with Zr, as Sc is known to provide additional strength through solid solution strengthening, precipitation hardening, and grain structure control [3,4]. The hardening potential of Sc was studied in the 1970s with a large part of the work conducted on binary Al-Sc [5,6] and ternary Al-Mg-Sc
⁎
systems [4,6–9]. Extensive reviews of the scientific literature on Sc in aluminium alloys were conducted in 1998 [3], 2005 [10] and more recently [11]. The strengthening possible from Sc in Al is very unique (and beneficial), and disproportional to the level of alloying when compared to other elements [12,13]. Despite the low solubility of Sc in Al, it achieves exceptional strength whilst retaining reasonable ductility. The addition of Zr modifies the Sc-dispersoids by both refining their size and increasing their thermal stability [14–16]. As such, the combined addition of Sc and Zr in aluminium results in alloys with exceptional strength and improved recrystallization resistance [3,14,17–21]. The Sc atoms were also reported to form ternary AlCuSc phases [22], if carefully designed heat treatment schedule to favour the formation of Al3Sc and θ′ was not utilised [23]. Despite these benefits, Sc has found almost no commercial applications so far because it was previously considered to be prohibitively expensive. However, with the projected price reduction of Sc in the next few years, there is a recent regain of interest for these alloys. The strength contributions from Sc and Zr additions are multi-fold: 1) solid solution strengthening, 2) precipitation of Al3(Sc,Zr) and 3) grain structure control. Sc atoms were recently reported to act as
Corresponding author. E-mail address: [email protected] (T. Dorin).
http://dx.doi.org/10.1016/j.msea.2017.09.032 Received 12 July 2017; Received in revised form 5 September 2017; Accepted 8 September 2017 Available online 09 September 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 707 (2017) 58–64
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heterogeneous nucleation sites for the θ′ phase and to then segregate at the θ′/matrix interface [24,25]. There is no available research on the effect of adding both Sc and Zr to an Al-Cu alloys. In this study, we examined the effect of alloying a binary Al-Cu alloy with different Sc and Zr contents. Four alloys were selected for the study as follows: Al-4 wt%Cu, Al-4 wt%Cu-0.2 wt%Sc, Al-4Cu-0.1 wt% Zr, Al-4 wt%Cu-0.1 wt%Sc-0.14 wt%Zr. The alloys were extruded, solution treated and aged. The effect of composition on precipitation and recrystallization was investigated using electron microscopy and atom probe tomography. The strength increment during ageing was monitored with tensile tests. It was found that the yield strength could be increased by up to 150 MPa with the suitable Zr and Sc content. An Al3Zr shell was found to form around the Al3Sc dispersoids. This paper is the first to experimentally observe the nucleation of θ′ precipitates on the core-shell dispersoids. The strengthening contributions are analysed and we find that the refinement of the θ′ precipitates accounts for the major part of the strength increase.
strain during the tests. At least 3 samples were tested under each condition to ensure reproducibility. 2.3. Microscopy The state of recrystallization was evaluated with electron backscattered diffraction (EBSD) measurements in a Jeol JSM 7800F scanning electron microscope (SEM) at 20 kV. The EBSD samples were sectioned and metallographically prepared using the usual methods. The final polishing step used was 5–10 min of polishing using colloidal silica. The precipitates distribution was characterised in a JEOL 2100F transmission electron microscope (TEM) equipped with energy dispersive spectroscopy (EDS). TEM samples were prepared by electropolishing in a solution of 33% nitric acid in methanol at −20 °C. Atom probe tomography (APT) was carried out on a LEAP 4000 h instrument at a pulse fraction of 200 kHz, a temperature of 60 K, and pulse fraction of 20%. Samples were prepared by electro-polishing in a micro-loop apparatus under an optical microscope in a solution of 5% nitric acid in ethanol at room temperature [27,28]. Electro-polished tips were given a final sharpen using Ga ions at 10 kV in a Quanta dual beam focused ion beam instrument.
2. Materials and experimental methods 2.1. Alloys preparation
3. Results
The alloys used for this study were cast using master alloys to ensure maximum homogeneity of the cast metals. Sc and Zr additions were made via standard master alloys compositions; Al-2 wt%Sc and Al-5 wt %Zr. It should be noted that Cu additions were made without the use of a master as it is known to dissolve easily into liquid Al [26]. Melting was conducted in a small induction furnace with a ~ 4 kg capacity for aluminium. The molten metal was kept at 710 °C for 15–20 min before being poured into a cylindrical steel mould. The cylindrical mould comprised a large steel base to provide directional solidification. The compositions of the alloys cast for this study can be seen in Table 1. The as-cast alloys were then homogenised to form Sc and Zr dispersoids and to dissolve the coarse intermetallic phases left over from the casting process. At least four homogenised billets were then extruded for each composition. The extrusion process was conducted on a specially designed module adapted to a traditional Instron apparatus. A small resistance furnace was used to pre-heat the billets prior to extrusion. The billet was then forced into a die at the selected temperature. A circular die of 5.3 mm diameter was used for all the extrusions conducted in this project. It should be noted no extrusion defects such as incipient melting or broken surface were observed in any of the resulting F-temper extrusions. The F-temper extrusions were then solution heat treated at 500 °C for 1 h before being stretched and aged at 160 °C for various times in order to form the hardening θ′ phase. The alloys were allowed to naturally age for 72 h prior to artificial ageing.
3.1. Strengthening kinetics In order to monitor the mechanical properties evolution during the ageing process at 160 °C, tensile tests were conducted on the four alloys at different artificial ageing times. Tensile specimens were aged for 1 h, 3 h, 10 h and 20 h at 160 °C. The tensile stress/ strain curves are displayed in Fig. 1(a)–(e) for respectively alloy 1 to alloy 4. A clear trend of simultaneous increase in strength and decrease in ductility with increasing alloying content can be observed. The yield stress, tensile stress and uniform elongation evolutions for all 4 alloys are reported in Fig. 2(a)–(c). The highest yield stress is reached after 10 h at 160 °C which is defined as the peak aged condition for all 4 alloys. The most significant increase in strength was observed in alloy 4 where the yield stress was 150 MPa higher than that of alloy 1. A drop in uniform elongation as a function of ageing time was reported for all 4 alloys. After 10 h at 160 °C, the uniform elongation ranged from 6% to 8%. 3.2. Recrystallization To evaluate microstructural differences, electron backscattered diffraction imaging was conducted on alloys 1, 2, 3 and 4 after the solution treatment of 1 h at 500 °C (see Fig. 3). Alloy 1 was observed to be fully recrystallised with a grain size of about 30 µm. Alloy 2 was also fully recrystallised but the grain size was refined with an average measured grain size of 20 µm. Alloys 3 and 4 displayed a strong extrusion texture with elongated grains. Whilst alloy 3 was partially recrystallised near the surface of the extrusion, alloy 4 was completely unrecrystallised. These observation are in agreement with previous studies that report a more efficient grain pinning from combined Sc and Zr additions as compared to the single elements.
2.2. Mechanical testing Tensile tests were carried out in an Instron 5967 equipment mounted with a 5 kN load cell. All tests were conducted at constant cross head velocity of 1 mm/min. For a 16 mm gauge length sample, this constant cross head velocity is equivalent to a strain rate of ~ 0.001/s. A non-contact video extensometer was used to measure the
3.3. Precipitation Table 1 Compositions of the four alloys under study.
Alloy Alloy Alloy Alloy
1 2 3 4
Al
Cu
Sc
Zr
Fe
Ti
Bal Bal Bal Bal
4 4 4 4
0 0 0.2 0.1
0 0.1 0 0.14
0.07 0.07 0.07 0.07
0.005 0.005 0.005 0.005
To examine the effects of Sc and Zr on the precipitation behaviour, two samples from alloy 1 and 4 in the peak aged condition were observed in transmission electron microscopy (TEM). The samples were observed in the <100> zone axis as two of the four θ′ variants can be observed edge-on in this orientation. It was found that the length and width of the θ′ plates were significantly decreased in the alloy that contained Sc and Zr (see Fig. 4). In 59
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Fig. 1. Tensile stress-strain curves for samples aged for 0 h, 1 h, 3 h, 10 h and 20 h at 160 °C for (a) Alloy 1 to (d) Alloy 4.
Fig. 6(c)).
alloy 1, the plates’ width was 197 nm on average and the thickness was 5–10 nm. In Alloy 4, the average plates’ width was 38 nm and thickness 2–5 nm. The θ′ plates size distributions was reported in the form of histograms, as shown in Fig. 4. The distributions both exhibit a lognormal shape. The presence of dispersoids was also observed, as shown in Fig. 5. The dispersoids were confirmed to contain both Sc and Zr and a coreshell type contrast was observed. In most instances, the dispersoids were observed to be in contact with a θ′ plate. This suggests that the dispersoids play a significant role in the nucleation and growth of θ′ which results in a refined distribution. In order to confirm the core-shell nature of the Sc and Zr containing dispersoids, alloy 4 was examined in atom probe tomography (APT), Fig. 6. Two dispersoids were present in the APT volume and the coreshell structure was clearly confirmed. The Sc rich core was observed to be about 5 nm diameter and the shell had a total diameter of about 10 nm. Fig. 6(b) reveals that the dispersoids are embedded in the centre of the θ′ precipitates which confirms that the dispersoids act as preferential nucleation sites for these precipitates. The composition evolution across the plates and dispersoids confirms that the dispersoids cores are Sc rich and that the shells contain both Sc and Zr (see
4. Discussion 4.1. Precipitation mechanisms Al-Cu alloys have been extensively studied and the main strengthening phase is known to be the metastable θ′ phase [29]. A traditional way to promote and refine these precipitates is by introducing additional nucleation sites which is traditionally performed by prestretching up to 6%. This strategy does result in an increase in strength but has the main disadvantage of decreasing the alloy's ductility and fatigue performance in the peak aged condition. Furthermore, in some processes, such as casting or forging, it is not possible to introduce cold deformation prior to ageing which decreases the potential hardening obtained from θ′ precipitation. Here, it was observed that dispersoids with a Sc-rich core and Zr-rich shell act as efficient nucleation sites for θ′ particles and could be used as an additional strategy to promote the formation of the main hardening phase with little or no cold deformation required. Additions of Sc and Zr are known to result in the formation of L12 dispersoids. These dispersoids traditionally form during solidification 60
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Fig. 2. Evolution of (a) yield stress (MPa), (b) ultimate tensile stress (MPa) and (c) uniform elongation (%) as a function of ageing time at 160 °C for the 4 alloys under investigation.
The θ′ precipitates observed in the peak aged temper for alloy 4 are clearly associated with the dispersoids. This suggests that the dispersoids act as preferential nucleation sites for these precipitates. The size of the θ′ precipitates was significantly different in the sample with
and high temperature homogenisation. In the present study, the alloys were homogenised prior to extrusion, hence, the dispersoids were already formed and stable. The combined additions of Sc and Zr were observed to result in precipitates with a Sc-rich core and Zr-rich shell.
Fig. 3. EBSD map at 100× magnification for respectively (a) Alloy 1 to (d) Alloy 4.
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Fig. 4. High magnification TEM micrographs showing Cu precipitates in (a) Alloy 1 and (b) Alloy 4. Note the difference in the scale bar. (c) and (d) Respective precipitates size histograms.
the θ′/matrix interface. We suggest that the Sc observed around the θ′ plates, in previous works [24,25], comes from the nucleation of θ′ on the Al3Sc dispersoids and then of a redistribution of the Sc atoms around the θ′ plates. In the present work, the presence of the Zr shell prevents Sc redistribution. Thus the Zr stabilizes the Al3Sc whilst keeping the refinement effect of the θ′ plates.
Sc and Zr as compared to the sample free from Sc and Zr. Indeed, the presence of Sc and Zr in alloy 4 refined the average diameter of the plates to 38 nm as compared to 197 nm with no Sc nor Zr in alloy 1. As a result, the dispersoids were found to help nucleate and stabilise small θ′ precipitates, preventing rapid coarsening of these precipitates. Contrarily to previous reports, there was no Sc segregation observed at
Fig. 5. TEM-EDS analysis of Alloy 4 (Al-Cu-Zr-Sc) aged at 160 °C for 10 h.
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Fig. 7. Strength contribution from the θ′ precipitates σ θ′ as a function of precipitate diameter as predicted from Eq. (2) and measured on alloys 1 and 4 from this study.
1
σ θ′ = 0.13MGα 2
(2)
where M = 3 is the Taylor factor, G = 28 GPa is the Al shear modulus, b = 0.286 nm is the magnitude of the Burgers vector and r0 = 0.572 nm is the inner dislocation cut-off radius. The evolution of σ θ′ as a function of precipitate diameter is plotted in Fig. 7. α = 7.4 was used as fitting parameter to get the best match with the empirical measurements. This value is smaller than what was reported elsewhere where α was found to be greater than 20 [30–32]. As a conclusion, the increase in strength contributions from the θ′ precipitates comes from the formation of much smaller precipitates in the presence of Sc and Zr. As such alloys 1 and 4 have average precipitates diameter of respectively 197 nm and 38 nm and this corresponds to a strength contribution from θ′ increasing from 50 MPa to 120 MPa.
Fig. 6. Atom probe tomography reconstruction of an alloy 4 sample aged at 160 °C for 10 h.
4.2. Strength contribution from θ′ The strength increment from the θ′ precipitates can be obtained by computing the strength increment during the final natural and artificial ageing. Indeed, for all the studied alloys, the θ′ precipitates form during the final ageing treatment, hence, the strength increment during ageing corresponds to the contribution to strength of the θ′ precipitates only. These contributions are summarised in Table 2. The alloys 1–4 have the exact same Cu content, hence we can hypothesize that the same volume fraction of θ′ precipitates should form. The major strengthening differences observed, hence only come from differences in precipitate size. The size of the θ′ precipitates can be described with two parameters, a diameter Dp and a thickness tp . Previous empirical and phase field modelling studies [30–32] have observed a proportionality relationship for these two parameters such that:
Dp = α × tp
0.079Dp 1 b [0.23 + 0.038α 2 Dp +0.007α ]ln ( ) r0 Dp
4.3. Additional strength contributions Apart from the refinement of θ′ precipitates, the most significant strengthening contributions appear to come from
(1)
1) precipitation hardening from the Al3(Sc,Zr) dispersoids 2) grain structure control
where α is a constant that depends on the alloy composition and heat treatment conditions. Zhu and Starke [33] developed a strengthening model for unshearable plate-like particles applicable to the θ′ precipitates. The variables in the initial model are the volume fraction fv , average diameter Dp and thickness tp of the precipitates. In the present study, the bulk composition in Cu is the same in the four studied alloys and if we assume that the stoichiometric volume fraction of fv = 0.051 is reached for all alloys after ageing to peak strength. As a result, Zhu and Starke's model can be computed as a function of the precipitate diameter Dp , and with one fitting parameter α , as follows:
To estimate the importance of these contributions, the tensile data of the as-quenched samples are used. As such, the strength of the asquenched alloy 1 (free from dispersoids) is used as a reference and is subtracted from the strength of the other samples as quenched. Indeed, all alloys contain the same amount of Cu so it can be considered that after the solution treatment, the only difference in the strength measured should come from the dispersoids and the resulting grain structure. The calculated strength increment are reported in Table 3. The observed strength increment is similar for alloys 2 and 3 with only Zr or only Sc which confirms that addition of Sc only does not provide
Table 2 Measured yield strength increment coming from the formation of θ′ precipitates.
Table 3 Measured yield strength increment coming from the L12 dispersoids and from grain size refinement for the 4 studied alloys.
Alloy
1
2
3
4
Alloy
1
2
3
4
σ θ′ (MPa)
50
74
76
120
σdisp (MPa)
/
17
20
40
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alloys used in this project. Deakin University’s Advanced Characterisation Facility is acknowledged for use of the Jeol JSM 7800F scanning electron microscope, Jeol 2100F transmission electron microscope and LEAP 4000 HR atom probe tomography instrument.
significant strength increase. However an extra 20 MPa is obtained by the addition of both Sc and Zr in alloy 4. The strength differences in the as-quenched condition is relative to the dispersoids’ volume fraction. Considering an Al3M stoichiometry and based on the alloying content from Table 1, the fraction of dispersoids should be similar in Alloy 3 and 4 which confirms that Al3(Sc, Zr) dispersoids are most potent hardener than the Al3Sc ones. Similarly, the stoichiometric fraction of dispersoids should be lower in Alloy 2 as compared to Alloy 3 and a similar strength contribution is observed in these alloys confirming again that the addition of Sc alone does not work efficiently in these alloys (Table 3). From this data, one realises that the most significant contribution to strength is by far the refinement of the θ′ precipitates (more dispersoids = more nucleation sites) rather than the direct hardening effects from the dispersoids or reduced grain size.
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5. Conclusions In the present study, the effects of adding Sc and Zr to an Al-4 wt %Cu alloy were investigated. Four alloys with different Sc/Zr ratio were analysed. The most interesting results can be summarised as follow:
• The yield strength was doubled for the alloy Al-4 wt%Cu-0.1 wt%Sc• • • • •
0.14 wt%Zr as compared to its binary counterpart. The peak strength was found to occur after 10 h of ageing at 160 ᴼC for all four alloys. Recrystallization was fully inhibited for the alloys containing both Sc and Zr. The Sc and Zr dispersoids were observed to form a core rich of Sc and a shell rich in Zr. The dispersoids helped the nucleation and stabilisation of θ′ precipitates during ageing. The θ′ precipitates were refined from an average diameter of 197 nm in alloy 1–38 nm in alloy 4. The contribution to yield strength from the θ′ precipitates was found to be up to 120 MPa. The contribution from the dispersoids + grain size was up to 40 MPa. Hence, the most important contribution to strength came from the role of the dispersoids in refining the θ′ precipitates.
Acknowledgements The authors would like to acknowledge Clean TeQ for providing inkind Al-Sc master alloys. Dave Gray is warmly thanked for casting the
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