Grain size

Normally, fine grain size is favourable for the mechanical properties; therefore, grain refiners are usually added in commercial alloys. However, since the resistivity is very sensitive to the composition, no grain refiner is used in the current study. Hence, the grain sizes of different alloys are distinctive due to the particular influence of each alloying element.

As a rough conclusion, the influence of Gd and Sn on the grain size is negligible and Al and Zn can refine the grain size. More specifically, in Mg-Al alloys the grain size decrease from

20 μm 20 μm

20 μm 20 μm

Segregation area

Segregation area

a)

d) c)

b)

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841 μm to 187 μm as the Al increases from 0.8 at. % to 8 at. % and in Mg-Zn it decreases from 312 μm to 166 μm with the increasing Zn contents from 0.8 at. % to 2.5 at. %. On the other hand, in Mg-Gd and Mg-Sn alloys grain sizes are approximately 1000 μm despite the changes of the solute content.

The grain refinement of as-cast alloys can be classified into two types:

1) Increasing the number density of nuclei during solidification;

2) Restricting the growth of grains.

The number density of nuclei can be increased by enhancing undercooling, or heterogeneous nucleation via introducing foreign particles and/or in situ synthesizing particles as effective sites. Restricting the growth of grains, on the other hand, can be achieved by adding alloying elements according to the concentration gradient mechanism [202]. The effect of solute elements on grain size can be assessed by the growth restriction factor (GRF) [203]. The GRF can be calculated using

𝐺𝑅𝐹 = 𝑚𝐶(𝑘 − 1) (6-1)

Where m is the slope of the liquid line (suggest it is a straight line), C is the concentration of the alloying element and k is the equilibrium distribution coefficient. The parameters for calculating GRFs of Al, Gd, Sn and Zn in Mg are calculated using the binary phase diagrams and listed in Table 6-2.

Table 6-2 Parameters for calculating GRFs of different alloying elements

Elements m k m(k-1)

Al -6.04 0.38 3.74

Gd -2.80 0.64 1.01

Sn -2.74 0.35 1.78

Zn -5.95 0.11 5.30

Zr [203] 6.90 6.55 32.89

Y [203] -3.40 0.50 1.70

A higher value of the GRF means more effective in restricting the growth of the grains and results in finer grain size. As in Table 6-2, when the concentration of each alloying elements is at the same level, then the GRF values of them are Zn>Al>Sn>Gd. This implies that Mg-Zn

71 alloys have the best effect in restricting the growth of the grains while Mg-Gd alloys have the worst effect, which is coincident to the results in Chapter 5.

6.1.2 Solution treated alloys 6.1.2.1 Cast alloys

The solution treatment for the cast alloys used to dissolve all the as-cast intermetallic phase into the matrix. Thus, the influence of different alloying elements and their contents on resistivity can be measured.

As in Fig. 5-12, after solution treatment, the as-cast intermetallic phases in 8Al and Mg-2.5Sn alloys have been dissolved into the matrix. Therefore, it can be concluded that the solution time is long enough for the Mg-Al and Mg-Sn series alloys to dissolve the as-cast intermetallic phases. Because Mg-8Al and Mg-2.5Sn alloys have the highest contents of the alloying elements and thus the highest amounts of intermetallic phases in Mg-Al and Mg-Sn series alloys. Consequently, they need the longest time to dissolve the intermetallic phases into the matrix compared to the alloys with lower contents. Despite the solute concentration being different, the solution time is same for each alloy series as indicated in Table 4-2. With the fact that the intermetallic phases in Mg-8Al and Mg-2.5Sn alloys are completely dissolved, it can be speculated that the intermetallic phases in Mg-Al and Mg-Sn series alloys have been dissolved.

In Fig. 5-12, some white rectangular phases are present in Mg-2.5Gd alloy and confirmed to be GdH2 by synchrotron radiation diffraction analysis as in Fig. 5-14. The hydrides have been found in many RE containing Mg alloys [204-208]. Vlček et al. [204] suggest that there are four ways in which the RE hydrides are formed:

1) they are formed during casting by the reaction of the RE element with hydrogen dissolved in raw materials;

2) they are developed as a result of hydrogen-induced decomposition of the as-cast intermetallic phases during solution treatment;

3) they are brought out due to the water quenching after solution treatment;

4) they are introduced during the sample preparation in the presence of water.

In fact, except for the third mechanism, all other mechanisms have been verified in Mg-Gd alloys by researchers [206, 208, 209]. In the current study, the as-cast alloys do not show obvious GdH2 peaks in the X-ray patterns and the samples are prepared without water.

72

Therefore, the GdH2 is formed mainly due to the hydrogen-induced decomposition of Mg5Gd during solution treatment.

Nevertheless, the RE hydrides are quite stable; they will not decompose up to 1000 °C [209].

Therefore, their amounts will not change during the current isothermal ageing temperature range (from 200 °C to 250 °C) and their influence on the resistivity is assumed as a constant.

Hence, it can be regarded as a systematic error and it will not affect the tendency of the resistivity change.

In the Mg-2.5Zn alloy, intermetallic phases are still present after solution treatment and they are considered as undissolved Mg7Zn3 phase. Usually, the existence of the as-cast intermetallic phases means an insufficient solution treatment duration. A longer solution time is needed to dissolve the intermetallic phases. However, things may be different in the current instance. As in Table 5-1, the concentration of Zn in Mg-2.5Zn alloy is 2.68 at. %, which is higher than the desired. According to the phase diagram in Fig. 6-1, the maximum solubility of Zn in Mg is 2.3 at. % at 340 °C. Therefore, the content of Zn is beyond its maximum solubility in Mg; in this case, a longer solution time will not help dissolve the Mg7Zn3 phase. A BSE graphic of the solution treated Mg-1.5Zn alloy with the same solution parameters is shown in Fig. 6-3. The BSE graphic shows that the intermetallic phase is completely dissolved into the matrix.

Therefore, the results indicate that the solution time is long enough to dissolve the as-cast Mg7Zn3 phase when the Zn content is less than the maximum solubility. The undissolved Mg7Zn3 phase in Mg-2.5Zn alloys is due to the excess Zn, not the solution time.

Fig. 6-3 Microstructure of solution treated Mg-1.5Zn alloy.

T4 treated Mg-1.5 Zn

100μm

73 The peaks shift in Fig. 5-13 is due to the different atomic radii of the solute elements. The atomic radius of Mg is 145 pm and the radii of Al, Zn and Gd are 118 pm, 142 pm and 233 pm respectively. The solute substitutes the Mg atoms in the matrix and thus changes the lattice parameters. Since the radii of Al and Zn are smaller than Mg, they will decrease the lattice parameters and, therefore, the interplanar distance of the matrix. On the other hand, Gd will increase the interplanar distance of the matrix. According to Bragg’s law, the 2θ value will change following the changes of the interplanar distance since the wavelength of the X-ray is a fixed value.

6.1.2.2 Extruded alloy

As mentioned in chapter 2, grain boundaries will increase resistivity. Since the grain sizes of the as-cast and solution treated alloys are quite different, it is necessary to clarify the influence of grain sizes on resistivity. Therefore, the extruded Mg-0.8Gd alloy is also investigated in the current study. The extruded Mg-0.8Gd alloy has an equiaxed and fine grain size, with different solution times, various grain sizes are obtained as shown in Table 5-12. The equiaxed grain size of the as-extruded alloy suggests that full recrystallization occurred during hot extrusion.

So the influence of the deformation-induced dislocations can be ignored. Hence, the only difference between these solution treated alloys is the grain size. They will be used to assess the influence of grain size on resistivity.

6.1.3 Aged alloys

As a general conclusion, the ageing temperature in the current study does not affect the precipitation sequence but accelerates the precipitation kinetics. In the Mg-Al and Mg-Sn alloy, the only observed precipitates are the Mg17Al12 and Mg2Sn phases, respectively. In the Mg-Gd alloys, the precipitates under the peaking condition are the Mg7Gd phase and under the over-aged condition, the Mg5Gd phase is observed. Although there is no direct observation of the 𝛽^′′ and 𝛽₁ phases due to the limited resolution of synchrotron diffraction. The observed results are coincident with the former studies. Therefore, the precipitation sequence of Mg-Gd alloy in the current study is thought to be SSSS → 𝛽^′′ Mg3Gd → 𝛽^′ Mg7Gd → 𝛽₁ Mg3Gd → 𝛽 Mg5Gd. In the Mg-Zn alloy, the Mg4Zn7 and MgZn2 phases are observed under peak-aged and over-aged conditions. According to Nie [5], the equilibrium precipitates in the over-aged condition is the MgZn phase. However, in the current study, the precipitates under the over-aged condition are still the MgZn2 phase; this may be attributed to the insufficient ageing time.

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Therefore, the precipitation sequence of Mg-Zn alloy is SSSS → 𝛽₁^′ Mg4Zn7 → 𝛽₂^′ MgZn2 in the current study.

Except for the precipitation kinetics, the ageing temperature also affects the growth of discontinuous precipitation in Mg-Al alloys. A higher ageing temperature is favoured for continuous precipitation, as shown in the microstructures that continuous precipitation occurs early for higher ageing temperatures. The formation and growth of continuous precipitates consume the supersaturated Al in the matrix, which will reduce the driving force for the growth of discontinuous precipitation. Also, small continuous precipitates act as the obstacles to pin the migration of grain boundaries hence inhibit the growth of discontinuous precipitation.

Therefore, a high ageing temperature can suppress the growth of discontinuous precipitation.