Your place for Aluminium Extrusion, Technology, Finishing and Applications

MICROSTRUCTURE AND PROPERTIES OF A NOVEL ZR-MICROALLOYED HIGH-STRENGTH AL MG SI ALLOY

F. Zupanič, T. Bončina, University of Maribor, Faculty of Mechanical Engineering, Slovenia
P. Cvahte, M. Steinacher, Impol Group, Slovenia
C. Gspan,Institute of Electron Microscopy and Nanoanalysis, Austria

Abstract                                                         

In this work, we investigated a modified Al-Mg-Si alloy (AA 6082). The main new alloying element was zirconium, however, also the additions of other elements were slightly modified. The alloy was characterised in the as-cast condition, after homogenization, extrusion, and T6 heat treatment using light microscopy, scanning and transmission electron microscopy with different microanalytical techniques, X-ray diffraction, differential scanning calorimetry, tensile, fatigue, and corrosion tests. In the modified AA 6082 alloy, the tensile strengths between 460 MPa and 500 MPa were achieved. The low-cycle fatigue strengths were comparable to similar alloys. The corrosion resistance of the alloy was slightly better in comparison to other copper-containing 6xxx alloys that are on the market (EN AW 6110A) but slightly worse than in 6xxx alloys without copper.

The excellent mechanical properties are based on the presence of Al₃Zr and -AlMnSi-dispersoids and nanosized ’-Mg₂Si and Q’-AlCuMgSi precipitates in the matrix.

Introduction

The manufacturing of green vehicles, which are characterised by lower fuel consumptions and CO₂ emissions, is a must for the automotive industry. This can be achieved through a reduction of the vehicle weight by replacing heavier steel parts with lighter, high-strength and, at the same time, tough aluminium alloys. Especially Al-Mg-Si alloys (series 6xxx) can be used. In this contribution, we investigated a novel Zr-microalloyed high-strength Al-Mg-Si alloy [1].

By development of a novel alloy, the complete chemical composition of the alloy was optimised [2]. The most important are zirconium and copper additions. Zirconium can decrease the size of crystal grains during solidification. However, the presence of zirconium in the alloy can reduce the efficiency of Al-Ti-B grain refiners. It is essential that zirconium form Al₃Zr dispersoids at higher temperatures. These dispersoid particles considerably decrease the tendency to crystal growth during heat treatment and hot plastic deformation. In this way, zirconium contributes to strengthening of the alloys by dispersion hardening and hardening by grain boundaries. Namely, the latter hardening is more important when the crystal grains are fine [3, 4].

The strengthening effect in Al-Mg-Si alloys is achieved through ' in '' precipitates [5]. These precipitates are metastable variants of the equilibrium Mg₂Si phase. They turn to Mg₂Si when the alloy is held prolonged at higher temperatures (between the ageing and solvus temperatures). One obtains a combination of 6xxx and 2xxx alloy by the addition of copper to Al-Mn-Si. In these alloys, the phase Q-AlCuMgSi dominates and not -Al₂Cu, which is predominant in 2xxx alloys. During ageing, Q’ precipitates are formed in the -Al, which strengthen the matrix in addition to '' precipitates [6]. The modification of Al-Mg-Si alloys by the additions of zirconium and copper enhances some strengthening mechanisms, leading to tensile strengths close to that of 7xxx alloys, but retaining much higher toughness.

The main aim of this study was to determine the microstructure of a novel Al-Mg-Si alloy after each step of the manufacturing process and to determine some mechanical properties and corrosion resistance.

Experimental

The alloy was prepared in an induction furnace (Junkers) at a temperature of 760 °C. The zirconium was added with the master alloy AlZr10. The melt was poured into a holding furnace for 90 min at temperatures between 730 in 740 °C. The billets with the diameter of 282 mm were semi-continuously cast using the AirSlip Casting Technology (Wegstaff). The addition of an AlTi5B1 grain refiner was 3 kg/t of the melt. The chemical composition of the alloy is given in Table 1. The billets were homogenised at different temperatures for a different duration and extruded into bars with a diameter of 57 mm. After extrusion, the bars were solution treated at temperatures between 530 °C and 550 °C up to 3 h and then exposed to artificial ageing at temperatures between 160 °C and 180 °C up to 12 h (T6 treatment).

Table 1: The chemical composition of the modified 6082 alloy.

The differential thermal analysis was carried out using Mettler Toledo 851. A sample was heated and cooled with the rate of 10 °C/min. The samples in the as-cast condition were metallographically examined using a light microscope Nikon EPIPHOT 300, and two scanning electron microscopes FEI Helios Nanolab 650 and FEI Sirion 400 NC. In SEMs, we also carried out microchemical analysis using EDS. X-ray diffraction was done at synchrotron Elettra (Trieste, Italy), by using X-rays with a wavelength of 0.099996 nm. A software Thermocalc and database TCAL5 was used for the simulation of solidification according to the Scheil model, and for calculating equilibrium phases as a function of temperature.

Lamellae for transmission electron microscopy (TEM) were prepared using a focused ion beam (FIB by FEI, Helios) for the as-cast and heat-treated specimens. High-resolution TEM (HRTEM), energy-filtered TEM (EFTEM), scanning TEM (STEM), and energy-dispersive X-ray spectroscopy (EDXS) were carried out in an FEI Titan80-300 image corrected electron microscope.

The tensile tests were done by Zwick Z250 according to ISO 6892-1 standard, and the fatigue tests by Instron 8802 according to EN ISO 26203-2. The samples were tensile-pressure loaded (R = -1) with frequency of 50 Hz at stresses of 200, 250, 300, 350, and 400 MPa.The corrosion tests were done according to PV1113 specification.

Results

XRD revealed clearly six phases in the as-cast condition: -Al, MgSi₂, tetragonal Al₃Zr, -AlMnSi, ZrSi₂ and -Si. The metallographic analysis further disclosed two minor phases Q-AlCuMgSi and Al₂Cu. Figure 1 shows a typical microstructure of the alloy in the as-cast condition. The equiaxed dendritic crystal grains of the aluminium solid solution -Al prevailed in the as-cast microstructure. The linear intercept length of crystal grains was around 235 m. Among the intermetallic phases, the volume fraction of the cubic -AlMnSi was the highest. EDS analysis confirmed that this phase also contained Fe and Cr. -AlMnSi was present in the interdendritic and intergranular spaces. The black Mg₂Si had a shape of Chinese script within a two-phase (-Al + Mg₂Si) structure. Zirconium was present in platelike particles of the tetragonal Al₃Zr and within the orthorhombic Si₂Zr. Si₂Zr crystallised on the Al₃Zr platelets and was not separated from Al₃Zr. The closer microstructure analysis revealed that -Si, Al₂Cu and Q-AlCuMgSi were present within the islands that formed via a multiphase reaction at the terminal stages of solidification (Fig. 1b). The Scheil solidification model predicted all discovered phases [7]. The solidification of such multi-component alloy is very complicated. DTA-analysis showed that it took place through several reactions (Fig. 2). The last melt solidified at approximately 540 °C. This temperature was slightly lower than the calculated solidus temperature of the alloy at 575 °C (Fig. 3).

Figure 1: The back-scattered electron micrographs of the alloy in the as-cast condition. a) A lower magnification showing the grain structure and the distribution of the main microstructural constituents, b) a detailed image of a multiphase island formed at the end of solidification.

Figure 2: A section of the heating DTA-curve between 500 °C and 750 °C (the heating rate 10 °C/min, argon) showing six peaks occurring during the melting of the alloy.

Figure 3: Calculated equilibrium phase fractions as a function of temperature. The calculated equilibrium solidus temperature T_s was 575 °C.

Fig. 3 shows the calculated equilibrium phases in this alloy as a function of temperature. It is clear that it is impossible to obtain a homogeneous solid solution by homogenisation. Several intermetallic compounds are stable just below the alloy’s equilibrium solidus temperature. The volume fraction of a-AlMnSi varies slightly with temperature. Thus this phase does not change significantly during homogenisation. Fig. 4 indicates only the change of shape of a-AlMnSi and undissolved Al₃Zr particle from AlZr10 master alloy.

On the other hand, the volume fraction of Mg₂Si decreased considerably. The remaining Mg₂Si did not have the shape of Chinese script but of globulites. During the optimal homogenisation treatment, the volume fraction of Mg₂Si decreased from around 1.4 % to 0.3 %. The dissolved Mg and Si are required for the precipitation of b’-precipitates during artificial ageing. Also b-Si, Al₂Cu and Q-AlCuMgSi that were present within the islands in the as-cast condition, dissolved completely. The TEM-investigation revealed that also Si₂Zr remained in the microstructure. It was attached to the micrometre-sized undissolved particle of a-AlMnSi and Al₃Zr. It was also found that a-AlMnSi and Al₃Zr dispersoids formed within the a-Al matrix during homogenisation (Fig. 4). They formed because of the supersaturation of the matrix with Zr and Mn after solidification. A detailed analysis showed that two types of Al₃Zr precipitates were present. The tetragonal precipitates had a shape of platelets, while the cubic Al₃Zr was present in the cuboidal shape. The TEM-micrograph (Fig. 5) also shows that dispersoids were distributed unevenly in the microstructure, which was caused by the inhomogeneous distribution of the alloying elements due to microsegregation during solidification. Al₃Zr dispersoids were typically in the proximity of a-AlMnSi dispersoids. They may be nucleated on a-AlMnSi. Therefore they did not actively contribute to the dispersion strengthening of the alloy. Still, zirconium probably had a significant role in decreasing the grain size of a-Al during the solidification and by the formation of a fine subgrain structure during hot pressing, and also contributes to the small grain size in the final T6-condition. 

Figure 4: The backscattered electron micrograph of the alloy after homogenization treatment. a) Lower magnification showing the grain structure and the distribution of the main microstructural constituents, b) higher magnification showing a change of shape of AlMnSi and Al₃Zr and a decrease of the volume fraction of Mg₂Si.

Figure 5: The transmission electron micrograph of the alloy after homogenization showing dispersoids within a crystal grain.

Figure 6 shows a TEM-micrograph of the alloy after hot pressing. The crystal grains are oriented in the pressing direction. In the areas with a high density of dispersoids, the grain sizes were 10 mm in the longitudinal and 0.5 mm in the radial direction. On the other hand, the grains were much larger in the areas with a lower density of dispersoids, indicating the importance of dispersoids in preventing grain growth.

Figure 6: The transmission electron micrograph of the alloy after extrusion showing the grain structure and distribution of dispersoids.

The final microstructure and properties were achieved by the T6 treatment, which consisted of the solution annealing and artificial ageing. During solution annealing, the crystal grains became coarser, and their shape changed from the elongated to the equiaxed. The linear size of grains was 20 mm. It was observed that coarsening of dispersoids took place (Fig. 7a). Nevertheless, they were still somewhat effective in preventing the grain growth. Elongated b’-precipitates (Mg₂Si) and rather equiaxed Q’-precipitates (AlCuMgSi) formed in the matrix during artificial ageing (Fig. 7b).

Figure 7: The transmission electron micrographs of the alloy after T6-treatment: a) the distribution of dispersoids in the crystal grains and b) an Al₃Zr dispersoid and Q’ and b’-precipitates.

The alloy was strengthened by small grains, three types of dispersoids (a-AlMnSi, tetragonal and cubic Al₃Zr) and two types of precipitates (b’-Mg₂Si and Q’-AlCuMgSi). The combination of these strengthening mechanisms resulted in rather a high yield and tensile strengths. Fig. 8 shows a typical tensile diagram. The yield strength was 440-460 MPa, the tensile strength 475-495 MPa, while elongation at fracture was 8-11 %. The fracture was ductile. The pores predominantly formed at large undissolved a-AlMnSi and Mg₂Si particles. Fig. 9 shows the results of low-cycle fatigue test. It shows that the modified and unmodified alloys possess very similar properties. 

Figure 8: A typical tensile diagram of the alloy Al-Mg-Si in the T6 temper. (R_p0,2 = 440 MPa, R_m = 490 MPa, A = 10 %).

Figure 9: The results of  the low-cycle fatigue tests of the initial and modified alloy.

 
Fig. 10 shows the result of the corrosion test. The modified alloy possessed worse corrosion resistance than the unmodified alloy because of the higher content of Cu. Nonetheless, it was better than with other 6xxx-alloys that contain copper.

Figure 10: The corrosion attack on the a) initial and b) modified AA 6082 alloy.

Conclusions

The results of the investigation showed that tensile strengths between 460 MPa and 500 MPa could be achieved by the optimal T6 heat treatment. The low-cycle fatigue strengths were comparable to the similar alloys. The corrosion resistance of the alloy was slightly better in comparison to other copper-containing 6xxx alloys that are on the market (EN AW 6110A) but slightly worse than in 6xxx alloys without copper.

The excellent mechanical properties are based on the presence of Al₃Zr and a-AlMnSi-dispersoids in the matrix that prevent grain growth during heat treatment and nanosized b’-Mg₂Si and Q’-AlCuMgSi precipitates that effectively inhibit the dislocation glide.

Acknowledgements

The work was carried out in the framework of the Slovenian smart specialisation, programme Materials and Technologies for New Applications, MARTINA, No. OP20.00369 and Infrastructure programme UM I0-0029 financed by the National Research Agency ARRS.

References

[1] H. J. McQueen and O. C. Celliers, Application of hot workability studies to extrusion processing .3. Physical and mechanical metallurgy of Al-Mg-Si and Al-Zn-Mg alloys, Can. Metall. Q. 36 (1997) 73-86.

[2] P. Cvahte, M. Jelen, D. Vukasin and M. Steinacher, A high-strength Al-Mg-Si aluminium alloy and its manufacturing process, edited by EPO (2017).

[3] Y. Himuro, K. Koyama and Y. Bekki, Precipitation behavior of zirconium compounds in Zr-bearing Al-Mg-Si alloy, in W. J. Poole, M. A. Wells and D. J. Lloyd, Aluminium Alloys 2006, Pts 1 and 2: Research through Innovation and Technology,  Trans Tech Publications Ltd, Zurich-Uetikon, 2006, pp. 501-506.

[4] L. Li, Y. Zhang, C. Esling, H. Jiang, Z. Zhao, Y. Zuo and J. Cui, Crystallographic features of the primary Al3Zr phase in as-cast Al-1.36wt% Zr alloy, Journal of Crystal Growth 316 (2011) 172-176.

[5] H. Farh, K. Djemmal, R. Guemini and F. Serradj, NUCLEATION OF DISPERSOIDS STUDY IN SOME Al-Mg-Si ALLOYS, Ann. Chim.-Sci. Mat. 35 (2010) 283-289.

[6] D. J. Chakrabarti and D. E. Laughlin, Phase relations and precipitation in Al-Mg-Si alloys with Cu additions, Prog. Mater. Sci. 49 (2004) 389-410.

[7] F. Zupanič, M. Steinacher and T. Bončina, The as-cast microstructure of a novel Al-Mg-Si alloy in 58th IFC Portorož 2018 - International Foundry Conference Portorož 2018 (Portorož, 2018).