Elsevier

Materials Characterization

Volume 154, August 2019, Pages 353-362
Materials Characterization

Micro-segregation and precipitates in as-solidified Al-Sc-Zr-(Mg)-(Si)-(Cu) alloys

https://doi.org/10.1016/j.matchar.2019.06.021Get rights and content

Highlights

  • Sc containing L12 dispersoids are found to form during solidification in the presence of Mg and Si. These dispersoids do not form in alloys free of Mg and Si in the same casting conditions.

  • The as-cast dispersoids are found to be rich in Si which substitute for ~25% of the Al in the L12 structure. Mg is found to partition at the α–Al/Al3Sc interface. We show that the Si content inside the dispersoids is a function of aging time rather than bulk alloy composition.

  • The presence of solidification micro-segregation in Al-Mg-Si-Sc-Zr alloys is quantified and modelled using a one dimensional Scheil model. The formation in Mgsingle bondSi precipitates is found to occur in the inter-dendritic regions. These precipitate are nucleated on the dispersoids. The Mg-to-Si ratio in these precipitates varies from ~1.5 to ~2 as a function of precipitate size.

  • Most of the Sc and Zr is found to supersaturate despite the presence of Si. This suggests that it may be possible to develop aging treatments that will precipitate Sc/Zr dispersoids from the as-cast condition with no need for prior solution treatment.

Abstract

In this work, we investigate the as-cast microstructure of three model alloys, an Al-Sc-Zr, an Al-Mg-Si-Sc-Zr and an Al-Mg-Si-Cu-Sc-Zr alloy. Transmission electron microscopy and atom probe tomography are used to investigate the precipitates present in the microstructure after casting. We find that higher Si content results in the formation of dispersoids upon casting. The as-cast dispersoids are found to be rich in Si with Si substituting for ~25% of the Al in the L12 structure. Variations in the local alloy composition and dispersoids size is explained by the presence of solidification micro-segregation which is fitted using a one-dimensional Scheil model. This segregation is found to be responsible for the formation of Mgsingle bondSi rod-like precipitates in the inter-dendritic regions. These precipitates are found to nucleate on the as-cast dispersoids. A closer look at the solid solution reveals that most of the Sc and Zr is supersaturated after casting even in the presence of Mg, Si and Cu. We discuss the impact of these results on the design of a suitable heat treatment for 6xxx-series alloys with Sc.

Introduction

The 6xxx-series aluminium alloys are extensively used in the automotive industry as they exhibit an attractive balance between strength, formability and corrosion resistance. In order to achieve higher strength, modern 6xxx-series alloys contain Cu up to 1 wt% and reach a yield strength in the vicinity of 400 MPa [1]. The strength increment in these alloys comes from the formation of the β″ and Q′ phases [2]. However, Cu-rich 6xxx alloys have a significantly reduced processability and corrosion resistance [3]. An opportunity to develop a high strength, formable alloy with appropriate corrosion resistance is with the addition of scandium.

Scandium is one of the most potent strengthener in aluminium alloys. Most of the strength contribution in these alloys comes from the formation of the Al3Sc dispersoids. The following strengthening mechanisms apply to Sc: 1) precipitation hardening, 2) recrystallization inhibiting, 3) assisting the nucleation of other strengthening phases [4]. As compared to other dispersoid phases, Al3Sc has a good electrochemical compatibility with aluminium and hence does not affect its corrosion resistance [5]. When added below its maximum solubility, scandium usually stays supersaturated after casting and the dispersoids are formed during a subsequent heat treatment [6]. The addition of Zr, together with Sc, leads to the formation of an Al3Sc-core Al3Zr-shell structure. The Al3Zr shell is stable at high temperature and hence prevents Sc from moving across the interface and allows retaining a fine distribution of dispersoids when exposed to high temperatures during processing (i.e. during solution treatment) [7]. The main challenge of using Sc in 6xxx-series lies in the heat treatment, which must be designed to allow for a fine precipitation of both Al3Sc and Mgsingle bondSi precipitates.

The research on Sc additions in Al-Mg-Si-(Cu) alloys is limited and the precipitation sequence in this system is still largely unknown. In Al-Si-Sc alloys, impurity levels of Si (below 0.2 wt%) were reported to enhance the precipitation kinetics of Al3Sc [[8], [9], [10], [11], [12]]. Higher levels of Si were reported to be detrimental to the mechanical properties [13]. For the case of 6xxx-series alloys with Sc, there is a lot of discrepancy in the available literature with some work reporting a positive strengthening effect [[14], [15], [16], [17], [18], [19], [20], [21]] while other report only a decrease in strength when adding Sc to an Al-Mg-Si alloy [[22], [23], [24], [25]]. The point of commonality between the studies reporting a positive effect is the use of a tailored Sc-specific thermo-mechanical treatment. Specifically, Rohklin et al. [[17], [18], [19], [20],26] avoided solution treating the Sc containing alloys prior to processing and Lityńska et al. [[14], [15], [16],21] utilise a multi-step heat treatment with an initial annealing stage at 300–350 °C to ensure the precipitation of Sc. Contrasting this, all of the studies that saw no benefit from Sc utilised a more standard processing route for 6xxx-series alloys, which includes an initial high temperature solution treatment and no heat treatment stage which triggers the precipitation of Sc (in the range 300–400 °C). This suggests that while Sc has potential to improve the strength of 6xxx-series alloys, a tailored heat treatment process must be utilised to ensure its effective use. To the best of the authors' knowledge, there is no detailed research work looking at understanding the full precipitation sequence in Al-Mg-Si-Cu alloys that contain Sc and Zr.

The first step to optimize the path for precipitation in Sc and Zr containing 6xxx-series alloys requires a detailed knowledge of the as-cast microstructure. The present work explores three model alloy compositions: an Al-Sc-Zr, an Al-Mg-Si-Sc-Zr and an Al-Mg-Si-Cu-Sc-Zr alloy. This work is a continuation of [27]. The microstructure is investigated using a combination of scanning and transmission electron microscopy (SEM and TEM) and atom probe tomography (APT). SEM imaging helps identifying the presence of local compositional variations attributed to solidification micro-segregations. Whilst the Al-Sc-Zr alloy is supersaturated in the as-cast condition, the presence of fine L12 dispersoids is observed in the Al-Mg-Si-Sc-Zr-(Cu) alloys. The composition of these dispersoids is quantitatively characterized and it is found that ~25% of the Al is substituted by Si in the Al3Sc structure; Mg, Zr and Cu are found to substitute for Sc. Mg and Cu are mainly found to partition at the α–Al/Al3Sc interface. By comparing with data from prior work, we find that the Si content within the Al3Sc is not linked with the bulk Si content of the alloy, but rather with the duration of isothermal aging. A one-dimensional Scheil model is used to model the solidification micro-segregations which also explain the different compositions of the three APT volumes. Mgsingle bondSi precipitates are found to form in the inter-dendritic solute-enriched regions. The Mg-to-Si ratio of these precipitates is found to be ~1.8. The impact of these different findings on the path to an optimal precipitate microstructure in a Sc-containing 6xxx-series alloys is further discussed.

Section snippets

Materials and experimental methods

The model alloys used in this study were prepared using 99.999% pure Al and the following master alloys: Al-25wt%Mg, Al-20wt%Si, Al-33wt%Cu, Al-2wt%Sc and Al-5wt%Zr. A small induction furnace was used for melting the alloys. The liquid metal was maintained at 710 °C for 30 min, under an argon atmosphere, before pouring into a cylindrical steel mold (80 mm in diameter x 200 mm tall). The cylindrical mold had a large steel base to provide directional solidification. The cooling rate was measured

Scanning electron microscopy

Scanning electron microscopy was first conducted on the Al-Mg-Si-Sc-Zr sample (see Fig. 1). The sample was observed perpendicular to the solidification direction. The presence of Fe-Mg-Si intermetallics was observed to occur at grain boundaries. Sc was found to segregate at a few intermetallics but no Zr segregation was reported at this scale. In order to assess the presence of solidification micro-segregation a long line EDX scan (2 h) was conducted with a small voltage to reduce the size of

Discussion

The three model alloys studied in this work were solidified under the same conditions. TEM was first used to show that the Sc and Zr remain in solid solution for the ternary Al-Sc-Zr alloy (Fig. 3a). Fine spherical dispersoids were observed in the as cast condition for both the Al-Mg-Si-Sc-Zr and the Al-Mg-Si-Cu-Sc-Zr alloys (Fig. 3b and c). The presence of Cu was not observed to have any noticeable impact on the as-cast Sc-containing dispersoids.

Despite the formation of these Sc containing

Conclusions

In the present study, the impact of adding Mg, Si and Cu on the solidification behavior of an Al-Sc-Zr alloy was investigated. A model alloy approach was used and the main conclusions are as follows:

  • The ternary Al-Sc-Zr alloy is supersaturated in the as-cast state. The presence of fine dispersoids is observed in the alloys containing Mg and Si. The nucleation of these as-cast Al3Sc dispersoids is assisted by Si.

  • The as-cast dispersoids are found to be rich in Si which substitute for up to ~25%

Data availability

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

Acknowledgement

Dr. Thomas Dorin is the recipient of an Australian Research Council Australian Discovery Early Career Award (project number DE190100614) funded by the Australian Government. The authors would like to acknowledge Clean TeQ for providing in-kind Alsingle bondSc master alloys. Dave Gray is warmly thanked for casting the alloys used in this project. Deakin University's Advanced Characterisation Facility is acknowledged for use of the SEM JSM 7800F, TEM-FEG JEOL 2100F and LEAP 5000 instruments.

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