List of Figures
Figure 1.1 Typical applications of hypereutecticAl-Sialloys. ............................................ 2 Figure 2.1 Fatigue property comparison between cast iron and aluminum engines ......... 4 Figure 2.2 a) Alusil® low pressure die casting (LPDC) for new Audi V6 and V8 SI linerless engine with only 5.5 mm land width; b) The surface bare-bore technology to expose Si phases; c) Sikasil® high pressure die casting (HPDC) with cast in BMI grey iron liners for Volvo in-line 5 cyl. diesel block ............................................................ 7 Figure 2.3 Gibbs free energy curves v.s. T under equilibrium cooling process for pure Aluminum calculated by the FactSage TM ................................................................. 11 Figure 2.4 a) The plot of free energy change ∆G v.s. nuclei size r; b) The homogeneous nucleation; c) The heterogeneous nucleation .......................................... 12 Figure 2.5 Wetting effects on varying barrier interfacial energies illustrated for heterogeneous nucleation mechanisms .......................................................................... 13 Figure 2.6 Growth morphologies during solidification process relates to varying heat flow gradients .................................................................................................................. 14 Figure 2.7 a) Free movement of dendrites at beginning of solidification via the Mass feeding mode; and developing to a solid network where only the Interdendritic feeding can take place .................................................................................................... 16 Figure 2.8 a) FactSage TM calculated binary Al-Si phase diagram under an
increases as laser powder increases. As reported by our previous works [ 2 , 14 ], fortifying the laser power enlarged the size of molten pool, and then the relative density of SLM processed sample increases. After the input energy is excess, the porosity rapidly increases given to the “keyhole” formation (see in Fig. 2 (a)). On the other hand, the Al-50Si alloy presents a slight decrement in relative density with excess (as shown in Fig. 2 (b)), which can be contributed to the high flowability of metal liquid in high silicon hypereutecticAl-Sialloys [ 18 ]. In details, Al-18Si and Al-50Si samples possess highest relative density about of 96.5% and 99.4% respectively. According to the work of Aminul Islam et al. [ 19 ], the porosity has a great impact on tribological behaviour and material removal during wear. Therefore, in this work, the samples with similar relative density of 96.5% for both Al-18Si and Al-50Si are chosen for avoiding the influence of porosity on tribological behaviour. The XRD patterns of SLM processed Al-50Si and Al-18Si alloys are shown in Fig. 3 . It can be seen that only Al and Si phases can be ob- served, which indicates the no oxidation and nitridation take place during SLM process. In details, the silicon changes from secondary to main phase, when silicon content increases from 18% to 50%.
Nowadays, hypereutecticAl-Sialloys are used extensively for automotive applications. For example, cast components that experience service conditions of high temperature and load, such as engine blocks, are now routinely produced using the Al-Sialloys [1-6]. It is understood that as-cast microstructure of the material is much dependent on melt processing in casting operations. Molten metal pouring temperature and cooling rate (hence, solidification rate), for example, determines size and distribution of primary Si particles, as well as crystallographic characteristics of the Al-Si eutectic [7-9]. All these in many ways determine in-service properties of the cast components.
In-situ thermal analysis and neutron diffraction techniques were used simultaneously to evaluate the kinetics of the non-equilibrium solidification process of Al-19%Si binary alloy. Feasibility studies concerning the application of neutron diffraction (source: NRU nuclear reactor, Chalk River, ON) for advanced solidification analysis were undertaken to explore its potential for high resolution phase analysis. Neutron diffraction patterns were collected in the stepwise mode during solidification between 740 and 400 o C. The variation of intensity of the diffraction peaks was analysed and compared to the results of a conventional cooling curve analysis. Neutron diffraction was capable of detecting nucleation of the Si phase (primary and eutectic), as well as the Al phase during Al-Si eutectic nucleation. Neutron diffraction reveals the presence of Si peaks about 23 o C above the liquidus temperature, as established by thermal analysis (i.e., 672 o C). This illustrates the potential of neutron diffraction for high resolution melt analysis at near-liquidus temperatures, required for advanced studies of grain refining, eutectic modification, etc. The solid-to-liquid volume fraction was determined based on the change of intensity of neutron diffraction peaks over the solidification interval. Overall, the volume determined was in good agreement with the results of the cooling curve thermal analysis. This study will help to better understand the solidification mechanism of Al-Sialloys used for various component casting applications.
Thermal analysis of the various 354 alloy melts was carried out to determine the sequence of reactions and phases formed during solidification under close-to-equilibrium cooling conditions. The main reactions observed comprised formation of the α-Al dendritic network at 598°C followed by precipitation of the Al-Si eutectic and post-eutectic β-Al 5 FeSi phase at 560°C; Mg 2 Si phase and transformation of the β-phase into π-Al 8 Mg 3 FeSi 6 phase at 540°C and 525°C; and lastly, precipitation of Al 2 Cu and Q-Al 5 Mg 8 Cu 2 Si 6 almost simultaneously at 498°C and 488°C. As a result of the low solidification rate of the thermal analysis castings, and a Zr content of 0.25 wt%, all Zr- containing alloys are located in the L + Al 3 Zr region of the Al-Zr phase diagram during the melting stage. Three main reactions are detected with the addition of Ni, i.e., the formation of AlFeNi, AlCuNi and AlSiNiZr phases. Larger sizes of AlFeNi and AlCuNi phase particles were observed in T alloy with its higher Ni content of 4 wt%, when compared to those seen in S alloy at 2% Ni content. Mn addition in Alloy U helps in reducing the detrimental effect of the β-iron phase by replacing it with the less-detrimental Chinese script α-Al 15 (Fe,Mn) 3 Si 2 phase and sludge particles. The Sc-intermetallic phases observed in this study appeared in two different forms: (Al,Ti)(Sc,Zr) and (Al,Si)(Sc,Zr,Ti). With the use of the multi-step solution treatment – involving higher solution temperatures and longer durations, an increased amount of incipient melting is expected to occur. Coarsening of the Si particles is also observed; with larger particles growing bigger at the expense of smaller ones. Primary Si particles are observed in the microstructure of the hypereutectic alloy L with its high Si content of 16 wt%.
As known well, due to its shorter process duration and lower capital investment requirements, T5 treatment is more cost effective than T6 temper. However, the downside of T5 temper is that hardness, yield strength and ultimate tensile strength of T5 heat treated castings are lower than those in T6 heat treated castings. Hence, optimization of heat treatment parameters is required to meet expected mechanical properties for new designed alloys. In order to do that, the response of the experimental alloys for various heat treatment applications should be carefully analyzed. In this study, an extensive heat treatment plan is designed and carried out for the high- pressure vacuum die casting alloys. The experimental alloys were exposed to T5 and T6 tempers and then, heat treatment parameters were determined according to hardness values of heat-treated alloys. As explained in experimental procedure (Chapter 3), T5 heat treatment temperatures were selected as 170°C and 210°C with various soaking times whereas T6 heat treatment temperatures were chosen as 500°C, 520°C and 540°C. For T6 tempers, solution heat treatment was carried out from 1 h to 24 hours. Then, it is followed by water quench and artificial aging at 170°C for 4 hours. 5.2.1 Hardness Evaluation of T5 Heat Treated Alloys
Figure 6.20. Micrograph of rosette from DTA sample with cooling rate at 5°C/min after holding at 595°C for 30 minutes.
From the morphology analysis of several samples, very fine eutectic silicon was observed within the rosette. In some of the rosettes, silicon plates also appeared encircling the rosette as seen in Figure 6.16. As stated earlier, the formation of rosettes in this case is more likely due to the liquid entrapment between dendrite arms. This finding is also in agreement with Terzi et al.  investigation on the dendrite coarsening mechanism and showed the liquid entrapment between dendrite arms in their in-situ tomography result. Attempts to see the 3D morphology of rosette from figure 6.20 was performed by deep etch technique. SEM micrograph showed the rosette not only as a small sphere, but deep etch sample indicated a form of cylinder shape with fine eutectic. Micrograph also revealed the presence of intermetallic phase within the rosette which is noticed from its plate shape (red arrow in Figure 6.21).
become semi-coherent during further coarsening . Despite the relatively low volume fraction of Al 3 Sc resulting from the low solubility of Sc even at high temperature, these
precipitates provide a significant strengthening effect . Al 3 Sc precipitates are
resistant to coarsening  giving rise to some good thermal stability of the microstructure up to 200°C. They are also very effective in stabilizing deformation substructure and in pinning grain boundaries during recrystallisation [21, 22]. This stabilization effect of deformation substructures is clearly exhibited in the low magnification TEM bright field image of the as delivered Al-Mg-Sc alloy of the present study (Fig. 6 (a) and (b)). These pictures show elongated low angle boundaries resulting from the cold rolling process (rolling direction is horizontal). The thickness of the domains is in a range of 0.1 to 0.5 μm with a typical length in a range of 1 to 5 μm. To exhibit the small coherent Al 3 Sc precipitates, the TEM foil was orientated in (002)Al
The properties of aluminum casting alloys can be improved through the appropriate control of several metallurgical factors involved in the production of these castings. Pure aluminum has low tensile strength but can easily be alloyed with other elements, for example, copper, zinc, magnesium, manganese, and silicon (i.e. duralumin), to make aluminum alloys which are the materials mostly used in industry. Only a few of them are used as major alloying elements in commercial aluminum-based alloys; others are used as supplements to alloying additions for the improvement of alloy properties and characteristics. The effects of these alloying additives on the properties of aluminum depend on the individual elements and the specific amounts added, as well as on their interaction with aluminum and with each other. While these additions are primarily used for strengthening, other elements are used to obtain specific microstructural characteristics which may include a finer grain size, higher critical recrystallization temperatures, or else to block the harmful effects of certain impurities. The alloying elements in aluminum alloys may be present in the form of solid solution, dispersoids, precipitates within the grain, or intermetallic compounds at the grain boundaries.
Two Al-Si foundry alloys, defined as 356-T7 and 319-T7, were prepared in the present work with commercially pure Al (99.7%), pure Mg (99.9%), Al-50%Si, Al-25%Mn, Al-50%Cu, Al-10%Sr and Al-5%Ti-1%B master alloys. The chemical compositions of experimental alloys are shown in Table 1 (all alloy composition in the present work is in wt. % unless indicated otherwise). After casting, T7 heat treatment was carried out by a 2-step solution treatment (500°C/4h+540°C/2h for 356-T7 and 495°C/4h+515°C/2h for 319-T7, respectively), followed by water quenching and artificial aging at 200 °C for 5 hours.
4.2. Application to cast aluminium alloys being investigated
In the section 3, it was shown that in the porosity-free alloy (alloy C), the crack initiation is principally related to persistent slip bands formation in the aluminium matrix, regardless of the loading mode. However, in the porosity-containing alloys (alloys A and B), fatigue cracks always initiate from casting porosity for both uniaxial and combined tension-torsion loads. These observations suggest that the probabilistic concept presented in the previous paragraph would be a reasonable choice to model the fatigue behaviour of the cast aluminium alloys under investigation. In this section, the choice of the criteria corresponding to each fatigue damage mechanism will be described as well as the procedure used to identify the model parameters. The predictions will be then compared with experimental fatigue strength data for the three alloys (see Tab.3).
économiquement significative en raison du prix plus bas du manganèse comparativement à celui du nickel.
La modification thermique des particules de silicium s'est révélée plus efficace avec des alliages modifiés au Sr plutôt qu'avec des homologues sans Sr. L'évolution des particules de silicium au cours des traitements en solution prolongée a suivi les mêmes tendances et séquences pour les alliages non modifiés de type 354 et 356 ainsi que pour les alliages de type 354 et 356 modifiés par Sr à différentes vitesses d'évolution. Le grossissement des particules de Si eutectique a été obtenu grâce à la coalescence des particules et aux mécanismes de mûrissement d'Ostwald; les deux mécanismes étaient actifs en même temps; Cependant, ils ont fonctionné de manière indépendante et additive. Les piqûres qui peuvent exister dans les particules de silicium peuvent être comprises comme l'impression (impression ou empreinte) laissée derrière la agglomération et la diffusion de petites particules avec / dans une (des) plus grande (s) particule (s).
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Cast aluminium alloys have been used to manufacture automo- tive components for many years due to their low density and excellent thermal conductivity. These components are subjected to cyclic mechanical loads which can cause fatigue damage and the functional failure of the structure. It has been demonstrated [1–6] that in high cycle fatigue (HCF), one of the principal param- eters controlling the fatigue strength of cast Al-Sialloys is the pres- ence of different microstructural heterogeneities which occur in the form of inclusions (Si particles and inter-metallics), casting defects (micro-shrinkage or degassing pores) and at the level of the aluminium matrix (often characterised by the DAS (Dendrite Arms Spacing) and/or SDAS (Secondary Dendrite Arms Spacing)
The aim of this paper is (a) to propose a flexible modelling framework to predict the fatigue strength as a function of defect size and (b) to apply this to cast Al-Sialloys, in
combination with the methodology proposed by Murakami  to characterize the casting defect size. The ultimate aim of this work is to propose simple analytical
To cite this version : Fernandez-Calvo, Ana Isabel and Niklas, Andrea and Lacaze, Jacques Comparison of thermal analysis and differential thermal analysis for evaluating solid fraction evolution during solidification of Al-Sialloys. (2010) Materials Science Forum, vol. 649. pp. 493-498. ISSN 1662-9752
To cite this version : Fernandez-Calvo, Ana Isabel and Niklas, Andrea and Lacaze,
Jacques Comparison of thermal analysis and differential thermal analysis for
evaluating solid fraction evolution during solidification of Al-Sialloys. (2010)
Materials Science Forum, vol. 649. pp. 493-498. ISSN 1662-9752