Hunter Sürekli Döküm Prosesi İle Üretilen Al-fe-mn-si Alaşımının Mikroyapısının Karekterizasyonu Ve Mekanik Özelliklerinin Belirlenmesi

Abstract

Bu çalışmada Hunter Sürekli DökümProsesi ile üretilen Al-Fe-Mn-Si yassı döküm alüminyum alaşımının 6 mm. ve 10 mm. kalınlıklarda dökümü yapılıp, mikro yapıları incelenmiş, her iki kalınlıktaki alaşımın mikro yapı farklılıkları ortaya konulmuştur. Malzemelerin döküm hallerinde ve farklı tav sıcaklıkları sonrasında yapılarındaki metaller arası bileşikler tespit edilmiş, uygun görülen bir homojen tav sonrası uygulanan üretim prosesi boyunca mekanik ve mikro yapısal değişimler takip edilmiştir. Alaşımın üretildiği yassı mamul alüminyum döküm teknolojisi ile ilgili olarak ve Al-Fe-Mn-Si sistemindeki değişik element konsantrasyonuna bağlı olarak bulunabilecek fazlar hakkında bilgiler verilmiştir. Her iki kalınlıkta malzemenin yüzeyde ve kesitte soğuma hızları hesaplanmıştır. Malzemelerin mikro yapılarının 6 mm. ve 10 mm. kalınlıklarda birbirine benzer olduğu görülmüştür. İncelenen alaşımın yüzeyinde ve kesitinde SEM incelemeleri yapılarak yapıda mevcut fazların kimyasal bileşimleri tespit edilmiştir. Yine bu malzemelerin yüzey ve kesitinde faz incelemeleri yapılmış, kübik ? fazının, döküm yapısında ve tavlı yapılarda kararlılığını koruduğu görülmüştür. Dolayısıyla alaşımın mekanik özelliklerinin kübik ? fazı tarafından direkt olarak etkilendiği anlaşılmıştır. Malzemelerde invers segregasyon tespit edilmiş, bunun alaşımın yüzeyinde kübik ? fazı oluşma şartlarının aleyhinde bir durum tespit ederken hegzagonal ? fazı oluşma şartları lehine bir durum yarattığı görülmüştür. 6 mm. kalınlıkta dökülen alaşım 560°C, 580°C ve 600°C sıcaklıklarda döküm sonrası 8 saat tavlanarak elde edilen mikro yapılarda, yapının ancak 580°C 8 saat tav sonrasında intedentritik örgünün tamamen kırıldığı dendridsel yapının değiştiği gözlenmiştir.

During the studies to set the mechanical properties of the specimens that is thicker than 200 micronn, Zwickl478 tensile testing machine, for less than 200 micron thick specimens MTS 200 tensile testing machine (thin aluminium) is used. For Spectral analysis ARL 3460 metal analyser, for Erichsen 125 deep suspension equipment is used, for the metalographic works Struers Lectropol-5 polishing machine, is made of Olypus PME 3 type optic microscope, SEM studies are with JOEL JSM 840, phase experiments are with Philips made material carried out. Aluminium flat alloy product covers a wide range of sectors including food, construction, medicine, otomotive and white goods. While these products usage area expands continuously, the quantity of these consumptions extends accordingly. As the worlds' aluminium reserves are limited and rapidly used, the necessity of the need to recycle this metal becomes vital in which from chemically scap aluminium more alloys which can be more useful to the production are tried to develop. During the course of this work, the examined Al-Fe-Mn-Si alloy seems to be less pure than the commercial pure aluminium alloys. Every year the world's aluminium industry produces a large amount of casting ingots to be rolled into thin sheets. The ingots usually contain iron, manganese and silicon as their main alloying elements. Depending on the exact alloy composition and the selected casting procedure, several different phases can result from the casting process. As the metal experiences the thermomechanical processing that follows the casting (e.g. homogenization, cold rolling, hot rolling and annealing) phase transformations can occur. Different phases usually have different properties and they also have different effects during nucleation and «crystallization. In order to produce good quality sheet it is necessary to know which phases the ingot contains at each stage in the production line. The phase transformation of Aİ6(Mn,Fe) into ? - Al(Mn,Fe)Si which is important in Al-Si-Fe-Mn alloy has been investigated by FURRER, Hausch et al. and Tromborg et al.[ 5 ] The aluminium sheets that are produced with the hunter dynamic casting have non-equilibrium structure. Non-equilibrium structures can be produced either by rapid cooling that suppresses invariant reactions or changes the nucleation of some equilibrium phase; or by heat treating to high temperature and quenching to retain the high temperature structure. In most aluminium alloys, complete retention of the high temperature structure is not possible and the movement toward equilibrium that takes place may produce new non-equilibrium phases and substantial changes of properties. In the Al - Fe - Mn - Si systems, (FeMn)Aİ6 is the first phase to form over a good part of the system where many of the commercial alloys are located. In many alloys (FeMn)Al6 then reacts peritectically with the liquid to form (FeMn)3Si2Ali5. Since the ratio iron/manganese in (FeMn)3Si2Alı5 can be higher then in (FeMn) Al« there is no need for iron rejection during the peritectic reaction and therefore the reaction can be fairly rapid. Thus in most commercial alloys the structure is the equilibrium one or close to it. However if the (FeMn)Al is in the form of massive crystals as is usual with very slow cooling the peritectic reaction may be incomplete and some free silicon may be present together with (FeMn)Al6. In high silicon alloys (FeMn)3Si2Alı5 may be primary and since its crystals tend to be limited by the (111) faces it appears as more or less well formed hexagons. Practically no manganese-bearing commercial alloys are in the fields where FeAl3, Fe2SiAlg or FeSiAU, are the first to crystallize. If manganese is added to aluminium-iron-silicon alloys the amount is close to the iron content to prevent those phases to form in appreciable amounts. If there are other elements (cobalt, chromium, nickel, etc.) which also combine with iron less manganese is needed. As in the ternary aluminium-iron-manganese alloys with cooling rates about l°C/sec, the FeAl« phase extends from one side to the other with lattice parameters somewhere between those of MnAİ6 and FeMs depending on the iron/manganese ratio. In the alloys in which (FeMn)Al6 is the first to form but that are close to the peritectic line (FeMn)Al6 -> (FeMn)3Si2 Ali5, fast cooling as can be obtained in continuous casting suppresses the nucleation of the (FeMn)Aİ6 phase and (FeMn)3Si2 Alı 5 may become the first manganese-bearing phase to crystallize. This results in a change of phase appearence: instead of having the sharp cornered structure of (FeMn)Ale the Chinese script is more rounded and smaller. Superheating of the melt which also tends to suppress nucleation of (FeMn)Aİ6 has the same effect. In the high silicon alloys fast cooling has no special effect it reduces the size of the phases of may result in the silicon assuming a more or less modified appearance but no new phases are formed or equilibrium one disappear. Splat cooling has the usual effect sharply increased solubility of iron, manganese, silicon but as to be expected for the same cooling rate the solubility of each element is less than in binary alloys. The little work that has been done on the heat treatment of the alloys has not shown any new feature.As in the aluminium-iron -manganese alloys the presence of iron tends to suppress the formation of MnAln and similarly to the aluminium-manganese-silicon alloys (FeMn)3Si2 AI15 precipitates together with (FeMn)Al6. In alloys in which silicon is present in amounts exceeding that needed for formation of (FeMn)3 Sİ2AI15 heat treatment may produce precipitation of silicon. Because of the limited solubility of silicon and the large difference in atomic diameter between aluminium and silicon the GP zones formed are very small and break away from the matrix after reaching the diameter of only some 10-20* 10"10 m. Thus the hardening due to silicon precipitation is very small and not worth commercial exploitation. [2] Al - Fe - Mn - Si alloys' mechanical properties were largely dependent on the substructure which was controlled by the homogeneus dispersion of intermetallic precipitate particles and their thermo-mechanical processing history. Small precipitates, distributed uniformly throughout the matrix, are desirable in order to obtain a small, stable cell structure and good mechanical properties. It is difficult to obtain a completely uniform distribution of the intermetallics in the as-cast matrix, since they precipitate as an interdentritic eutectic during solidification. The spacing of the particles, which depends on the dendrite arm spacing, determines the minimum cell size obtainable by thermal - mechanical processing. Their size has an effect on the alloy's response to mechanical working and resultant properties, inasmuch as large precipitates can decrease the recrystallization temperature by acting as nucleation centers. Small precipitates can increase the recrystallization temperature by pinning dislocation tangles and subgrain boundaries. Solution heat treatment cannot be used for precipitate control since the intermetallics are are insoluble. Consequently, control of the precipitate-matrix morphology is obtained by deformation and / or thermo-mechanical treatment. Deformation of the material causes the precipitates to be broken and strung out, thus increasing their numbers and decreasing their spacing. However the precipitates break up only if they are brittle and if sufficient deformation is applied prior to annealing so that high internal stresses, which exceed the precipitate's fracture strength, can be built up. Refinement of the final structure is, therefore, obtained by optimizing casting, working and homogenization conditions. In addition to affecting the dispersion of precipitates and substructure, all deformation processing produces an alingment of grain into preferred crystallographic orientations. This texture is essentially controlled by the crystal structure of the metal or alloys; the shape change that it undergoes; the quantity of alloying additions; the temperature of working; and the amount and type of reduction imposed. The development of particular textures may lead to a resistance to yielding, i. e. texture hardening. In this research, studies are carried out on justifying the characteristics of the hunter dynamic casting and flatly produced casted Al-Fe-Mn-Si alloy. First The microstructure cross section of material and surface between 6 mm. and 10 mm. thick casted alloys are examined and the microstructure difference between the two different thicknesses is determinated. In regard to the coolness speed when calculating the speed coolness of 6 mm. and 10 mm. thick alloys surface and its crossection, difference in mikrostructure is seen. Later at the end of the homogenization process for 8 hours at 560 °C, 580 °C and 600 °C on the structure of the alloys and casting structure of alloys, the phase and phase transformations are determined. Also SEM experiments have been carried out to give support to the phase work and to see chemical compositions in the structure of the existing alloy phase section. The material that is annealed accoding to the ideal anneal heat to determine the alloy, is rolled in different thickness and during the rolling process differences in mikrostructure and mechanical characteristics is determinated. In both the thicknesses not homogenaus, thin and coarse dentric structure is found together in the alloys structure.This structure dominates both the crossections and the surface of the material. The reason for this is because of the speed of cooling differences in the thin and coarse dentritic structure area. The dendritic structure is the coarser in the areas where it is slowly cooled. The size difference between thin and coarser dentric structure is seen more on the surface of the material than the crossection. Segregation is more dominant in the structure of the whole material all through the crossection. The cooling rate of 10 mm. thickness alloys, during casting: Fine dendride regions on the surface of alloy: min. 830 K/sn. Max. 4184 K/sn. Coarse dendride regions on the surface of alloy: min. 30 K/sn. Max. 121 K/sn. Fine dendride regions on the cross section of alloy: min. 1078 K/sn. Max. 2803 K/sn. Coarse dendride regions on the cross section of alloy: min. 55 K/sn. Max. 73 K/sn. The cooling rate of 6 mm. thickness alloys during casting: Coarse dendride regions on the surface of alloy: min. 34 K/sn. Max. 73 K/sn. Fine dendride regions on the surface of alloy: min. 2353 K/sn. Max. 3845 K/sn. Coarse dendride regions on the cross section of alloy: min. 32 K/sn. Max. 175 K/sn. Fine dendride regions on the surface of alloy: min. 1843 K/sn. Max. 3375 K/sn In the phase examining of the 10 mm. thick casting on the surface of the material, hegzagonal ? and Al6Fe compaund is found. But in the cross section of the material in addition to the hegzagonal ? and Al6Fe compaund there is also cubic a phase is found. Casting structure, together with hegzagonal ? and Al6Fe compaund, weak cubic a phase is found on the surface of the 6 mm. thick alloy. Also hegzagonal ?, Al6Fe and cubic a is determinated on the crossection of the 6 mm. thick alloy. The existing phase in the surface and crossection of the 6 mm. thick specimens subject to homogenization for 8 hour at 560 °C, 580 °C and 600 °C is found to be cubic a phase. With this result, its shows that the dominant phase in the microstructure of the homogenizated alloys is cubic a phase. In regard to this, cubic a phase directly effects the mechanic properties of the materials. Cubic a phase is a phase that is seen more in fast cooling conditions. With these results gained this approach seems to be clashing. Nevertheless with these materials invers. Segregation existence explains why the cubic a phase on the surface is formed very litter or not formed at all. The coolness speed of the fast cooled surface reduces with the liquid floatation towards the surface. Therefore the conditions of forming the cubic a phase is from far off. The existence of invers segration increases the formation chance of the hegzagonal a phase. It shows that at the end of applying homogenization at 560° C to the 6 mm thick material the interdentritic network in alloy structure is not broken complately.Dentritic structure in the alloys structure preserves itself in some areas, this kind of structure is not suitable for a beginning structure which is needed the produce the material that has the reguired mechanical properties. The amount of the existing phases for above mentioned structures can not be calculated because of the technical restrictions.Calculating the amount of these phases with a different study would enable us to understand better the connections between the materials mechanical properties and the existing phases. Other than that existing phases anneal heat and in connection to its duration the variations could be seen. After the homogenization treatment of the 6 mm thick material at 560° C for 8 hours, it is subject to another homogenization treatment at 580° C for 8 hours. Along with this treatment dentritic structure is lost,the interdentritic network totally broken, the homogeneous and eguazi morfology in the structure is dissolved, the obtained structure after the treatment of homogenization at 580°C for 8 hours, is suitable as a starting structure which is in possession of the reguired mechanical properties. Casted at 6 mm thickness and treated with homogenization at 580° C is rolled down to 30 micro thickness. Some micro structure variations are established during the steps of rolling.As the thickness of the interdentritic phase material gets thinner, it gets broken and the size reduces. While the 0.2 % proof stress and tensile strength of the materials increase, % alongation of the materials decrease.