Harnessing Oxides in Liquid Metals and Alloys

Transcription

1 Harnessing Oxides in Liquid Metals and Alloys H-T Li, Y Wang, M Xia, Y Zuo and Z Fan BCAST, Brunel University, Uxbridge, Middlesex, UB8 3PH, UK Abstract Oxides are inevitably present in liquid metals and alloys. They hamper solidification processes, degrade mechanical performance of final cast products and cause excessive metal loss. To mitigate such problems, expensive melt cleaning processes are often used to reduce/eliminate oxides prior to solidification processing. In this paper we demonstrate that oxides in alloy melts can be harnessed by intensive melt shearing, a melt conditioning technique developed by BCAST. We found that oxide films are liquid-like films comprising fine oxide particles densely populated in a liquid matrix. It is shown that intensive melt shearing can disperse such oxide films into individual oxide particles distributed uniformly in the alloy melt. Dispersing oxide particles in this way not only mitigates their harmful effects on solidification processing and mechanical properties, but also provides effective nucleation sites for grain refinement. After a brief review on oxidation of liquid Al- and Mg-alloys, we discuss the effects of intensive melt shearing on the nature of oxides, provide evidence of nucleation on oxide and exemplify intensive melt shearing as an enabling technology for solidification control. 1. INTRODUCTION Due to their high affinity for oxygen at high temperature, liquid metals and alloys oxidize naturally when they are exposed to oxygen-containing atmospheres. This means that liquid metals and alloys (referred to as alloy melts hereafter) inevitably contain oxides as inclusions. Such oxides present some difficult challenges during solidification processing of metallic materials [1]. Firstly, oxides exist in alloy melts as thin oxide films, which reduce the fluidity of the melt and cause blockage during die filling; secondly, the entrained oxide films often exist in the alloy melts as bifilms, which become cracks in the solidified materials and degrade mechanical properties; thirdly, severe oxidation decreases materials yield, and therefore increased component cost. Consequently, oxide inclusions in alloy melts are usually treated as undesirable substance in the foundry. The conventional wisdom is either to prevent/reduce oxidation by melting metallic materials in a protective atmosphere or in a vacuum, or cleaning the alloy melts by either chemical or physical means or both [2]. Both melt protection and melt cleaning can only be effective to some extent and are expensive operations. However, recent scientific understanding and technological development indicate that it is possible to harness oxides in alloy melts 93

2 94 Solidification Science and Technology by intensive melt shearing [3,4]. It has been demonstrated that intensive melt shearing prior to solidification processing not only mitigates the aforementioned harmful effects but also makes oxides useful for grain refinement during solidification processing [3] {Al 2 O 3 }/<γ-al 2 O 3 > G b (a) G, Jm h a c γ t γ s -1.0 γ i h a, nm G, Jm G b G t {Al 2 O 3 }/<γ-al 2 O 3 > h a c γ s γ i (b) h a, nm Figure 1. Bulk energy ( G b ), interfacial energy ( γ i ), surface energy ( γ s ) and the total Gibbs energy ( G t = G b + γ s + γ i ), as functions of thickness (h a ) of the amorphous oxide {oxide} and the corresponding crystalline oxide <oxide> on the surface of liquid metals at 700 C. (a) The {Al 2 O 3 }/<Al 2 O 3 > system; (b) the {MgO}/<MgO> system. The critical thickness h a c is determined by the point at G t = 0. h a c is about 1.3 nm for the {Al 2 O 3 }/<Al 2 O 3 > system and 0.15 nm for the {MgO}/<MgO> system at 700 C (as indicated by the arrow).

3 The John Hunt International Symposium 95 In this paper we provide an overview of our work on harnessing oxides in alloy melts for grain refinement. We will confine our discussions to oxides in Al- and Mg-alloys, although the same approaches can be applicable to other alloy melts. After a brief review of oxidation of Al- and Mg-alloy melts in Section 2, we present the physical nature of oxides in Al- and Mg-alloy melts and their dispersion by intensive melt shearing in Section 3, grain refinement by oxides in Section 4, harnessing oxide in casting processes in Section 5 and a brief summary in Section OXIDATION OF Al- AND Mg- ALLOY MELTS Aluminium oxide has a large free energy of formation, and oxygen has very low solubility in liquid aluminium. This means liquid Al oxidizes readily at the surface when it is exposed to an atmosphere containing oxygen and/or water moisture [5]. The oxidation process starts with the formation of a thin amorphous layer on the melt surface. Based on the thermodynamic model for oxide overgrowth on a solid metal surface developed recently by Mittemeijer and co-workers [6], Men and Fan [7] have developed a thermodynamic model to analyze the stability of amorphous oxide on liquid metals. For the Al/Al 2 O 3 systems, the thermodynamic model predicts that the positive bulk Gibbs energy difference between amorphous and crystalline oxides can be compensated for up to a critical thickness by the negative energy difference of surface and interfacial terms, resulting in a thermodynamically stabile amorphous oxide layer below a critical layer thickness of 1.3 nm (see Figure 1a). The stable amorphous oxide layer is dense and continuous on the melt surface. Consequently, it not only separates the liquid Al from the oxygen containing atmosphere, but also hampers the diffusion of chemical species involved in the oxidation reaction since there is no fast diffusion path available in the amorphous structure [8]. However, beyond the critical thickness this amorphous layer becomes thermodynamically unstable and transforms into the stable crystalline oxide. It is found that after an incubation time of 5-10 minutes at 750 C, the amorphous layer transforms to crystalline alumina by means of nucleation and growth [9]. This crystalline form may be γ-al 2 O 3 or η-al 2 O 3 or others, depending both on alloying elements in the melt and the oxidation conditions [9,10]. The crystalline alumina nucleates and grows rapidly at the melt/oxide surface. It was reported that after the formation of the crystalline γ-al 2 O 3 or η-al 2 O 3 film the alumina will be subjected to further transformation to the α-al 2 O 3 crystalline form with increasing time and temperature. This transformation from γ-al 2 O 3 to α-al 2 O 3 is very sluggish in nature [11,12]. Currently there is no consensus on the time and temperature for this transformation in the literature due to its sluggish nature, diverse alloy systems and wide range of oxidation conditions. The complete transformation from γ-al 2 O 3 to α-al 2 O 3, however, has been reported to be achieved at 750 C after an incubation time of about 5 h in commercial purity (CP) aluminium (Al-0.15Fe-0.08Si, all composition are in wt.%, unless stated otherwise) [5]. However, in the industrial grade 319 alloy (Al-6Si-3Cu), the formation of α-al 2 O 3 was reported to occur at 850 C and was proposed to be related to the higher content of impurity elements in this alloy [12]. Mg is an important alloying element in Al-alloys, such as in AA5xxx and AA7xxx wrought Al-alloys. In Al-alloys containing Mg, the oxidation tendency increases sharply

4 96 Solidification Science and Technology with an increase of Mg content [5]. Depending on the Mg content, the oxidation reaction path possibly starts with the formation of amorphous MgO, MgAl 2 O 4, or Al 2 O 3, which then transforms to crystalline MgO, MgAl 2 O 4, or γ-al 2 O 3 films, respectively. Due to breakaway oxidation, the main oxides formed on the melt surface may be quite different when other kinetic factors are taken into account such as the melt temperature, holding time and turbulence on the melt surface. Haginoya and Fukusako [10] studied the oxidation behaviour of molten Al-Mg (2-12%) alloys at the temperature range between 650 C and 900 C for 3 h in dry air. It was identifiedd by quantitative X-ray diffraction analysis that MgO was produced at early stages of oxidation, and its amount increased temporarily and then decreased gradually, while the amount of MgAl 2 O 4 increased gradually with the decreasing amount of MgO. This is understandable. Mg is an effective surfactant to molten Al, and a high concentration of Mg at the melt surface allows for initial oxidation to form MgO. In addition, since MgAl 2 O 4 is thermodynamically more stable than MgO, MgO will give way to MgAl 2 O 4 with increasing reaction time. The oxidation reactions can be described by two consecutive reactions: Mg + 1/2O 2 MgO followed by MgO + 2Al + 3/2O 2 MgAl 2 O 4 [10]. Molten Mg oxidises readily in an atmosphere containing oxygen, often leading to complete burning on the melt surface if the melt is exposed to air and no precautionary measures are taken. In contrast to the case for liquid Al, the bulk Gibbs energy difference is relatively large for the Mg/MgO system, and cannot be compensated for by the surface and interfacial energy difference (see Figure 1b) [7]. This means that an amorphous MgO layer is thermodynamically unstable and crystalline MgO particles form directly from the melt. In order to process molten Mg safely and efficiently, the melt is normally protected either by exclusion of oxygen (e.g. covering with salt flux or inert gas) or by changing the nature of the surface oxide with a cover gas, to slow the oxidation rate down to an acceptable level. The compositions of the surface film of the alloy melt varies with the composition of the alloy and cover gas [13-17]. For molten AZ91D alloy, the surface oxide films are mainly composed of MgO and some MgAl 2 O 4 [17]. Figure 2. SEM micrographs showing the typical morphology of naturally occurring γ-al 2 O 3 oxide collected by pressurised melt filtration from commercially pure aluminium (CP Al) melt at 700 C. γ-al 2 O 3 oxide exists in the melt usually as (a) oxide films, and sometimes as (b) oxide bifilms.

5 The John Hunt International Symposium 97 (a) (b) 100µm 10µm (c) (d) 10µm 0.5µm Figure 3. OM (a) and SEM (b,c,d) micrographs showing the typical morphology of naturally occurring MgAl 2 O 4 oxide in different Al-alloys collected by pressurised melt filtration at 700 C. (a) MgAl 2 O 4 films in LM24; (b) MgAl 2 O 4 films in LM24 showing discrete MgAl 2 O 4 particles; (c) crumpled MgAl 2 O 4 films in LM24; (d) naturally dispersed MgAl 2 O 4 particles in Al-5Mg alloy. 3. PHYSICAL NATURE OF OXIDES IN Al- AND Mg- ALLOY MELTS 3.1. Morphology of Oxides in Alloy Melts Study of oxides in liquid metals is not an easy task since oxides in solidified materials are not easy to find and are difficult to identify. A pressurised melt filtration technique was used to collect oxides from the melt to facilitate microscopy studies. Figure 2a shows the collected oxide in CP Al at 700 C, which has been identified as γ-al 2 O 3 by X-ray diffraction [18]. γ-al 2 O 3 exists in CP Al as continuous films at low magnification. A closer examination reveals that γ-al 2 O 3 films consist of platelets and some time form bifilms entrained in the melt, as shown in Figure 2b. It is important to realise that γ-al 2 O 3 films are not completely dense and continuous and that γ-al 2 O 3 platelets are held together by liquid metal. This explains why γ-al 2 O 3 films can be so flexible in alloy melts. Oxide in liquid Al-alloy melts can be different from that in CP Al. The oxide in liquid LM24 and binary Al-Mg alloys has been identified as MgAl 2 O 4. Under the optical microscope with low magnification, MgAl 2 O 4 films appear as flexible dark lines (Figure 3a), while SEM examination with high magnification shows that MgAl 2 O 4 films consist of

6 98 Solidification Science and Technology discrete MgAl 2 O 4 particles held together by a liquid matrix (Figure 3b). With excessive oxidation or with increased scrap contents, MgAl 2 O 4 films can crumple forming an MgAl 2 O 4 particle cluster (Figure 3c). However, with increased Mg contents in Al-Mg alloys, MgAl 2 O 4 particles can be naturally dispersed, and are more likely to be uniformly distributed in the alloy melt. For instance, MgAl 2 O 4 collected by melt filtration from liquid Al-5Mg alloy is discrete individual platelets (Figure 3d). This natural dispersion (a) 18µm (b) 5µm (c) 7µm Figure 4. SEM micrographs showing the typical morphology of naturally occurring MgO oxide in AZ91D alloy melt collected by pressurised melt filtration. (a) Oxide skins formed on solid ingots before melting; (b) young oxide films formed during melt handling; (c) old oxide films formed at the melt surface in the furnace during melting.

7 The John Hunt International Symposium 99 may be attributed to decreased liquid/mgal 2 O 4 interfacial energy due to preferential segregation of Mg at such interfaces. The decreased interfacial energy leads to reduced capillary force between the MgAl 2 O 4 particles, so that even a mild shear flow, such as those created by melt handling and pouring, will be adequate to separate the MgAl 2 O 4 particles in the alloy melt. Oxides in Mg-alloy melts are dominantly MgO due to the significantly larger free energy of formation compared with those of other oxides. MgO particles are readily present in Mg-alloy melt through chemical reaction with oxygen during heating, melting and melt handling processes even though the melt is usually maintained in a protective atmosphere (Figure 4). Our experimental results show that there are three different origins for MgO present in AZ91D alloy melt, namely, oxide skins, young oxide films and old oxide films [3]. Oxide skins brought into alloy melts by solid ingots are usually µm in thickness and contain a high volume fraction of fine MgO particles, and therefore have a straight plate-like morphology (Figure 4a). The young oxide films in Mg-alloy melts originate from instantaneous oxidation of the fresh melt surface and the entrainment of such thin films in alloy melts. They exhibit flexible worm-like thin liquid-like films filled with a high volume fraction of nanometre-sized MgO particles and with a thickness of usually a few micrometres (Figure 4b). The old oxide film formed on the melt surface in the melting furnace is a result of prolonged oxidation under the protective atmosphere in the melting furnace (Figure 4c). In comparison with the young oxide films, old oxide films contain larger oxide particles (1-2 µm in size), and may have different chemical compositions depending on the chemical nature of the protective gas used Morphology of Oxides in Intensively Sheared Alloy Melts In recent years, BCAST developed a high shear technique for conditioning alloy melts prior to solidification processing [3,4]. Intensive melt shearing with either a twin screw mechanism [3] or a rotor-stator high shear device [4] can subject the alloy melt to an intensive forced convection with high shear rate and high degree of local turbulence. The shear rate achievable by the high shear technique is orders magnitude higher than that provided by propeller based mixers, electromagnetic stirring and other existing physical fields. Recent research has demonstrated that naturally occurring oxides and other inclusions present in alloy melts are amenable to physical manipulation by intensive melt shearing [3]. The local shear force provided by the intensive melt shearing is high enough to overcome the capillary force, so that oxide particles in oxide films can be freed from constrain of the capillary force. Consequently, oxide films are dispersed into individual particles, which are then distributed uniformly in the alloy melt. As shown in Figure 5, after intensive melt shearing, γ-al 2 O 3 films are converted into submicron-sized platelets in CP Al (Figure 5a), MgAl 2 O 4 films are dispersed into micron- or submicron-sized MgAl 2 O 4 particles in Mg-containing Al-alloy melts (Figure 5b&c), and MgO films are changed to discrete nanometre-sized MgO particles in Mg-alloy melts (Figure 5d). It should be pointed out that the clustering of oxide particles in Figure 5 is a result of pressurised melt filtration. It is reasonable to believe that oxide particles in the intensively sheared alloy melt are distributed uniformly throughout the entire volume of the melt.

8 100 Solidification Science and Technology (a) (b) 1µm 8µm (c) (d) 2µm 0.3µm Figure 5. SEM micrographs showing the typical morphology of naturally occurring oxides in alloy melts collected by pressurised melt filtration after intensive melt shearing. (a) Platelet γ-al 2 O 3 in sheared CP Al; (b) MgAl 2 O 4 in sheared Al-5Mg melt; (c) MgAl 2 O 4 in sheared LM24 melt; (d) MgO in sheared AZ91D melt. Intensive melt shearing disperses the oxide films into discrete particles Benefits of Dispersing Oxide Films As an enabling technology, dispersing oxide films by intensive melt shearing can deliver a number of beneficial effects to solidification processing of metallic materials. Firstly, dispersing oxide films facilitates casting processes, since it increases the fluidity of alloy melts at a given melt temperature, and avoids potential blockage by large oxide films during die filling. This means high quality castings can be produced at a lower melt superheat, resulting in an increased casting efficiency, lower energy consumption and longer die life. Secondly, dispersing oxide films can improve the quality of cast products by reducing the chances of the existence of large oxide films and inclusions in the castings, which is one of the major factors responsible for inferior mechanical performance of cast components compared with wrought products. It has been demonstrated that intensive melt shearing prior to casting can significantly improve the ductility of cast components [18]. Thirdly, dispersing oxide films can provide several orders of magnitude higher number density of potential nucleating particles to enhance heterogeneous nucleation for grain refinement. This will be further explained in the next section of this paper. Finally, dispersing oxide films allows an increased content of scrap

9 The John Hunt International Symposium 101 metals to be used in the melting furnace without degradation of the quality of cast components. Increased scrap metal content usually leads to excessive oxides and other inclusions in the alloy melt, causing processing difficulties and inferior casting quality. Intensive melt shearing can mitigate such problems by reducing/eliminating the harmful effects of oxides and inclusions without the deployment of expensive chemical treatment to eliminate the oxides and other inclusions themselves Average grain size, (µm) (a) Casting temperature, ( o C) 600 Average grain size, (µm) (b) Shearing temperature, ( o C) Figure 6. Average grain size of CP Al as a function of processing temperature. (a) Without intensive melt shearing; (b) with intensive melt shearing. Grain size was assessed on samples cast in the standard TP-1 mould with a fixed cooling rate of 3.5 K/s.

10 102 Solidification Science and Technology 4. GRAIN REFINEMENT OF Al- AND Mg-ALLOYS BY OXIDES 4.1. Grain Refinement by Intensive Melt Shearing As discussed previously, intensive melt shearing results in an effective dispersion of oxide films into discrete oxide particles and a uniform distribution of the dispersed oxide particles in the alloy melt, giving rise to a significant increase in number density of oxide particles as potential nucleation sites. Theoretically, this may lead to a substantial grain refinement if the oxide particles are potent for heterogeneous nucleation during solidification. Figure 6a shows the average grain size of CP Al assessed by the standard TP-1 test as a function of casting temperature (or melt superheat) without intensive melt shearing [18]. The grain size is at the millimetre level and increases with increasing melt superheat and then tends to decrease slightly with further increase of melt superheat. However, application of intensive melt shearing to CP Al leads to significant grain refinement, showing a complex grain size dependence on melt superheat (Figure 6b). The grain size increases initially with the increase in melt superheat, decreases sharply at 740 C from 500 µm to about 150 µm and then increases at a slower pace with further increase of melt superheat Average grain size, (µm) Non-Sheared Sheared wt.% Mg Figure 7. Average grain size as a function of Mg content in Al-Mg alloys with and without shearing. Grain size was assessed on samples cast in the standard TP-1 mould with a fixed cooling rate of 3.5 K/s.

11 The John Hunt International Symposium Average grain size, (µm) Non-sheared Sheared 800rpm 45s Processing temperature, ( o C) Figure 8. Average grain size of AZ91D as a function of processing temperature with and without intensive melt shearing. Grain size was assessed on samples cast in the standard TP-1 mould with a fixed cooling rate of 3.5 K/s. In addition, it is found experimentally that the effect of intensive melt shearing on grain refinement is also dependent on alloy composition [19]. Figure 7 presents the grain size of binary Al-Mg alloys solidified in the TP-1 mould with and without intensive melt shearing as a function of the Mg content. Without intensive melt shearing, the grain size decreases sharply with the increase of Mg content in dilute Al-Mg alloys (Mg <1%), and then levels off at around 200µm with further increase of the Mg content. It is interesting to note that intensive melt shearing of Al-Mg alloys can only be effective in alloy melts with low Mg content, and has no effect on grain size of the concentrated Al-Mg alloys (Mg >1%). Furthermore, intensive melt shearing is found to be extremely effective for grain refining Mg-alloys [3,4,20]. Figure 8 compares the grain size of AZ91D alloys solidified in the TP-1 mould with and without intensive melt shearing at different melt superheat. Without intensive melt shearing, the grain size increases with increasing melt superheat following the well-known S-curve (Figure 8). Compared with Al-alloys, superheat has a stronger effect on the grain size of Mg-alloys. The grain size is around 200 µm at melt temperatures close to the alloy liquidus (600 C), and is increased to around 700 µm at 650 C. However, such superheat dependence of grain size is significantly suppressed by application of intensive melt shearing, resulting in a consistently small grain size, only varying between 90 and 190 µm for the same range of superheat. These results are significant since, so far, there is no effective grain refiner available for Al-containing Mg-alloys [21].

12 104 Solidification Science and Technology 4.2. Potency of Oxides as Nucleation Substrates The experimental results presented in the previous section have clearly demonstrated that intensive melt shearing prior to solidification processing can result in significant grain refinement in both Al- and Mg-alloys without the addition of any chemical grain refiner. In this section we discuss the mechanisms of grain refinement by intensive melt shearing. As discussed previously, intensive melt shearing results in a significant increase in number density of oxide particles dispersed in the alloy melt. If such dispersed oxide particles are potent for heterogeneous nucleation, intensive melt shearing would lead to significant grain refinement. Therefore, the key factor to be assessed is the potency of oxide for nucleation in Al- and Mg-alloys. The potency of a solid substrate in a liquid metal can be defined as the degree of lattice matching across the interface between the substrate and the solid phase to be nucleated [22]. The lattice misfit f can be used as a quantitative measure of the potency for heterogeneous nucleation, and is defined as f = (d S d N )/d S, where d S and d N are the atomic spacing along a close packed direction on a close packed plane of the solid and the nucleating substrate, respectively. The calculated lattice misfit with solid Al at 660 C is 3.38% for γ-al 2 O 3, 0.48% for α-al 2 O 3, and 1.41% for MgAl 2 O 4, indicating that those oxides are highly potent for nucleation of α-al in consideration of the corresponding lattice misfit being 4.22% for TiB 2 and 0.09% for Al 3 Ti [22]. Although the calculated lattice misfit with the α-mg is 8.01% for MgO, the experimentally measured lattice misfit in AZ91D alloy is considerably less (5.46%) [3]. It is important to point out that it is the lattice misfit of the solid/substrate interface at the moment of nucleation that determines the potency of the nucleating system, not the theoretically calculated misfits for pure metals. Adsorption of solute elements at the liquid/substrate interface is a thermodynamically favourable process if such segregation lowers the interfacial energy [23]. As discussed elsewhere [22], the segregated solute elements at the interface may increase or decrease the potency of the substrate depending on whether they reduce or increase the lattice misfit. In addition, for enhancing heterogeneous nucleation to achieve effective grain refinement, the nucleating particles not only need to be potent, but also need to have an adequate particle number density, a suitable particle size and a narrow size distribution [24,25]. With such understanding, we can explain the experimental results presented in the previous sections. In CP Al, the oxide exists in liquid Al in the form of γ-al 2 O 3 films (Figure 2a). Although γ-al 2 O 3 is potent for nucleation of α-al, there is not a sufficient number density of γ-al 2 O 3 particles as nucleation sites, resulting in a coarse grain structure in CP Al (Figure 6a). With intensive melt shearing, γ-al 2 O 3 films are dispersed into submicron-sized γ-al 2 O 3 platelets (Figure 5a). Consequently, the number density of γ-al 2 O 3 particles is increased by several orders of magnitude, leading to a significant grain refinement (Figure 6b). In addition, intensive melt shearing enhances the kinetic condition for mass transport in the melt and hence accelerates any chemical reactions and phase transformations that occur in the system. Therefore, it is reasonable to assume that intensive melt shearing accelerates the transition from γ-al 2 O 3 to α-al 2 O 3 in liquid CP Al. Based on the experimental results in Figure 6b, we suggest that the equilibrium temperature for the γ-al 2 O 3 to α-al 2 O 3 transition is C. It is well-known that γ-al 2 O 3 to α-al 2 O 3 transition is accompanied a 24% volume reduction and fragmentation.

13 The John Hunt International Symposium 105 Hence, there will be a significant increase in particle number density after the transition. Since α-al 2 O 3 is more potent than γ-al 2 O 3 it is expected that there will be a large reduction of grain size after the γ- to α-al 2 O 3 transition due to the increased nucleation potency and particle number density, as shown in Figure 6b. Similarly, grain refinement by intensive shearing of dilute Al-Mg alloys (Mg <1%) and Mg-alloys can be explained by increased number density of potential nucleating particles. However, in the concentrated Al-Mg alloys (Mg >1%), it will be impossible for intensive melt shearing to produce further grain refinement since MgAl 2 O 4 particles in such alloy melts are already dispersed naturally as discussed previously in Section Evidence of Heterogeneous Nucleation on Oxides Nucleation on a potent substrate will result in a coherent or semi-coherent solid/substrate interface. Therefore, the most direct evidence of heterogeneous nucleation on potent substrates is the existence of an orientation relationship between the solid and the substrate obtained by electron diffraction, or an atomic image of a coherent or semi-coherent solid/substrate interface obtained by high resolution transmission electron microscopy (HRTEM). However, it has been proven to be very difficult to carry out such studies. It is well accepted that the number efficiency for heterogeneous nucleating particles is very low, being 0.1-1% [24]. This means that the majority of the added particles do not participate in the nucleation process, and that the probability for finding a nucleating particle in the TEM sample with suitable orientation for electron diffraction or HRTEM study is extremely low. Nevertheless, we have been successful to a certain extent at finding some evidence for nucleation on oxide, which is presented below. Figure 9 is a HRTEM micrograph of the MgO/α-Mg interface in AZ91D alloy indicating that the MgO/α-Mg interface is semi-coherent. This confirms that MgO particles can act as sites for heterogeneous nucleation during solidification of Mg-alloys. The derived orientation relationship between MgO and α-mg is [3]: ( 0001)[21 10] α Mg //(111)[011 ] MgO (1) The experimentally determined lattice misfit with this orientation is 5.46% [3], which is much smaller than the theoretically calculated lattice misfit of 8.01% for pure Mg presumably due to the segregation of solute elements at the MgO/α-Mg interface. In Al-Mg alloys, TEM examination has confirmed that the MgAl 2 O 4 particles have {111} planes as their naturally exposed surface [19]. Selected area diffraction patterns have determined the following cube-on-cube orientation relationship between MgAl 2 O 4 and α-al: ( 111)[110] MgAl O //(111)[110 ] 4 α Al 2 (2) This is very much expected since both MgAl 2 O 4 and α-al have the same crystal structure and closely matched atomic spacings along the close packed directions on the close packed planes. The calculated lattice misfit at 660 C is 1.41% along the [110] direction on the (111) plane. It is clear that the lattice misfit between MgAl 2 O 4 and α-al is quite

14 106 Solidification Science and Technology (1 1 1) MgO [0 1-1] ( ) Mg [ ] 2 nm Figure 9. High resolution TEM micrograph of the interface between MgO and the α-mg matrix in AZ91D alloy. The α-mg matrix (bottom) is viewed along [1-210] direction and MgO (top) is along [01-1] direction. MgAl O 2 4 [0 1 1] (1-1 1) α -Al [0 1 1] (1-1 1) 2 nm Figure 10. High resolution TEM micrograph showing a faceted MgAl 2 O 4 particle and its interface with the α-al matrix. Both MgAl 2 O 4 and the α-al matrix are viewed along [011] direction. The {111} crystal planes for the two phases deviated by about 18 from each other.

15 The John Hunt International Symposium 107 small compared with that for the Al/TiB 2 system ( 4.22%). Figure 10 is an HRTEM micrograph of the MgAl 2 O 4 /α-al interface in LM24 suggesting that the MgAl 2 O 4 /α-al interface is semi-coherent. It is noticed that the {111} crystal planes of MgAl 2 O 4 and α-al deviated by about 18 from each other. This may be attributed to the severe segregation of solute and impurity elements at the MgAl 2 O 4 /α-al interface. 5. HARNESSING OXIDE FOR GRAIN REFINEMENT IN CASTINGS Intensive melt shearing has been used as an enabling technology for harnessing naturally occurring oxides in alloy melts in various casting processes for both grain refinement and improved casting quality. Here are a few examples of such applications. The high shear device has been implemented into the direct chill (DC) casting process to form the melt conditioned DC (MC-DC) casting process (Figure 11a) [20]. Figure 11b shows the microstructure of an MC-DC cast 80mm AZ91 billet. The microstructure dramatically changes from a coarse dendritic microstructure without intensive melt shearing (bottom of Figure 11b) to a fine and uniform equiaxed microstructure with the high shear device being switched on (top of Figure 11b). This microstructural change has been attributed to enhanced heterogeneous nucleation, increased nuclei survival rate and potential dendrite arm fragmentation [20]. Similarly, intensive melt shearing has also been applied in the twin roll casting (TRC) process to form the MC-TRC process [26], which can provide Al- and Mg-alloy strip with fine and uniform microstructure without centreline segregation. (a) (b) 500µm Figure 11. (a) Schematic illustration of the melt conditioned direct chill (MC-DC) casting process; (b) Microstructure of the DC cast AZ91 alloy billet showing the dramatic change from a coarse columnar grain structure solidified without intensive melt shearing (bottom) to a fine and uniform equiaxed microstructure solidified with intensive melt shearing (top).

16 108 Solidification fication Science and Technology Figure 12. Microstructures of sand cast LM25 components with a 24mm thickness with and without intensive melt shearing. (a) Conventional sand casting (SC); (b) melt conditioned sand casting (MC-SC). (a) (b) 10mm (c) 10mm 10mm (d) 10mm Figure 13. Macrographs in cross-section section (a,b) and vertical section (c,d) showing the grain structure of CP Al solidified in the standard TP-1 mould without (a,c) c) and with (b,d) the addition of α-al2o3 based grain refiner.

17 The John Hunt International Symposium 109 Intensive melt shearing has been applied to shape casting processes. Sand casting is a slow cooling process, often producing an extremely coarse microstructure. Grain refinement is thus very desirable to improve the mechanical performance. Intensive melt shearing has been applied to sand casting (SC) to form the MC-SC process. Figure 12 compares the microstructures of sand cast LM25 with and without intensive melt shearing, suggesting that intensive melt shearing is very effective for grain refinement in sand cast components. Similar results have also been demonstrated in other shape casting processes, such as high pressure die casting (HPDC) and gravity die casting processes [18]. The fact that oxides are often potent substrates for heterogeneous nucleation of α-al and α-mg has prompted us to develop oxide based grain refiners using the intensive melt shearing technique to disperse synthetic oxide particles in master alloys. For example, synthetic α-al 2 O 3 particles with suitable size and size distribution have been dispersed in Al-alloys to form an α-al 2 O 3 containing master alloy for grain refining Al-alloys [27]. Figure 13 demonstrates the grain refining effect of α-al 2 O 3 based grain refiner on CP Al. Without addition of the grain refiner, CP Al has a coarse columnar grain structure (Figure 13a,c). However, addition of the α-al 2 O 3 based grain refiner results in a fine and uniform equiaxed grain structure (Figure 13b,d). 6. SUMMARY Due to their high affinity for oxygen Al- and Mg-alloy melts oxide readily in atmospheres containing oxygen. We found that the oxides formed naturally in Al- and Mg-alloy melts are in the form of continuous liquid-like oxide films consisting of densely populated oxide particles in a liquid matrix. Such oxide films are amenable to physical manipulation by intensive melt shearing to form dispersed oxide particles uniformly distributed in the alloy melt, resulting in a few orders of magnitude increase in number density. Theoretical analysis and experimental results have confirmed that most of the oxide particles in the sheared alloy melts are potent substrates for heterogeneous nucleation, have adequate number density, and suitable size and size distribution for grain refinement. Experimental evidence has been provided for nucleation on oxides by high resolution transmission electron microscopy and for significant grain refinement of Al- and Mg-alloys by TP-1 tests. It is demonstrated that oxides in Al- and Mg-alloy melts can be harnessed to mitigate their harmful effects and be deployed for grain refinement during a variety of casting processes. ACKNOWLEDGEMENTS The EPSRC is gratefully acknowledged for providing financial support. REFERENCES [1] Campbell J. Castings, 2 nd edn. Oxford: Butterworth-Heinemann; [2] Neff DV. In: ASM Handbook, vol. 15 Casting. ASM; [3] Fan Z, Wang Y, Xia M, Arumuganathar S. Acta Mater 2009;57:4891. [4] Fan Z, Zuo YB, Jiang B. Mater Sci Forum 2011;690:141.

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