Otomotiv sektörü için Al-Si-Fe-X alaşımlarının geliştirilmesi

Abstract

Otomotiv sektöründe büyük önem kazanan Al-Si-Fe-X alaşım sisteminin yapı özelliklerine hızlı katılaşma ve püskürtme biriktirimin etkilerini araştırmak ve Türk otomotiv sanayine katkıda bulunmak amacı ile, geleneksel kokil kalıba döküm yöntemi ile elde edilen alaşım sistemi; hızlı katılaşma melt-spinning yöntemi, atomizasyon ve püskürtme biriktirimi yöntemleri ile üretilerek döküm, katılaşma, ısıl işlem ve sertlik özellikleri incelenmiştir. Hızlı katılaşma M.S. yöntemi ile üretilmiş Al-Si-Fe alaşımı şeritlerde ortalama soğuma hızı yaklaşık 106-107oC/sn. değerindedir. Gaz atomizasyonu ile üretilen 500um. üstü Al-Si-Fe alaşımları toz numunelerin ortalama soğuma hızı yaklaşık 102oC/sn.'dir. 500um.-100um.aralığındaki tozların soğuma hızının da 102- 104oC/sn. olduğu tahmin edilmiştir. Püskürtme biriktirimi (Osprey) yöntemi ile üretilmiş Al-Si-Fe alaşımları ön-şekil numunelerin ortalama soğuma hızı 103- 104°C/sn. aralığındadır. Şerit numunelerde bakır tekerle temasta olmayan kenar yüzeyindeki yapının, bakır tekerle temasta olan kenar yüzeyindeki yapıdan daha kaba olduğu tespit edilmiştir. 500um. üstü Al-Si-Fe alaşımı toz numunelerin partikül şeklinin küresele yakın olduğu, 90um. altı toz numunelerin partikül şeklinin gözyaşı damlası şeklinde olduğu belirlenmiştir. Geleneksel kokil kalıba döküm, hızlı katılaşma M.S. yöntemi, atomizasyon ve püskürtme biriktirimi Al-Si-Fe alaşımlarının sertlik değerlerinin alaşımlardaki artan silisyum içeriği ile artmıştır. Hızlı katılaşmış Al-XSi-3.3Fe alaşımları şerit numuneleri geleneksel kokil kalıba döküm numunelerinden ortalama yaklaşık 3 kat daha büyük sertlik göstermiştir. Hızlı katılaşmış Al-XSi-3.3Fe-l.7Cu-0.5Mg şerit numuneleri geleneksel kokil kalıba döküm numunelerinden ortalama yaklaşık 2 kat daha büyük sertlik göstermiştir. Doğal yaşlandırma ısıl işlemi uygulanmış Al-XSi- 3.3Fe-1.7Cu-0.5Mg alaşımları kokil kalıba döküm numunelerinde silisyum ağırlıkça %8.4 olduğunda maksimum sertlik 140Hv. olup, alaşımdaki artan silisyum içeriğiyle ağırlıkça %10.4 olduğunda maksimum sertlik değeri 160Hv. ve %20 olduğunda 164Hv. değerine çıkmıştır. Doğal yaşlandırılmış püskürtme biriktirimi Al-8.4Si-3.3Fe-l.7Cu-0.5Mg ön-şekil numunesinde 120 saat sonraki maksimum sertlik değeri 128Hv. iken, alaşımdaki silisyum içeriğinin artışıyla Al-20Si-3.3Fe-l.7Cu-0.5Mg ön-şekil numunesinde 120 saat sonraki maksimum sertlik değeri 1 40Hv. olmuştur

CTTIVyfA/T AT>V U KJ lTJLXTXJ-».JLlk. JL Development of AI-Si-Fe-X Alloys For Automotive Industry Rapid Solidification Processing (RSP) clearly combines the main features of the overall field of materials science and engineering-processing, structure/property relationships. Rapid solidification processing was practiced industrially over 150 years ago in the production of lead pellets. Fast cooling undoubtedly resulted in refined dendritic microstructures, even possibly with some degree of supersaturation and fine-scale precipitation. Another important rapid solidification process was, production of soldering wire, patented in 1871 by E.M. Lang. The 1960s brought intensified interest in rapid cooling as an approach to finer dendritic structures and to compositional uniformity through smaller diffusional distances. This led the way to inert-gas atomization and ribbon-making processes. In addition, Pol Duwez' innovative splat-quenching technique provided access to increased degrees of supercooling, higher levels of supersaturation, and new metastable phases including metallic glasses. Three industrial developments during the 1970s ushered in the present era of rapid solidification processing. The Crucible Steel Company entered the tool- steel market with nitrogen-gas atomized high speed steel powders, consolidated into bulk stock by hot isostatic pressing. The Allied Corporation began to produce metallic glasses in ribbon and strip form by melt spinning. Pratt and Whitney Aircraft (Florida) developed the centrifugal atomizing of superalloys, combined with rapid quenching of the molten droplets in high-velocity helium gas, thus generating rapidly solidified powders under well-controlled conditions. The most recent general review of rapid solidification of metals was that by Jones, published in 1984. Specific reviews have been published for magnesium alloys in 1989, for titanium alloys in 1990, and for aluminum alloys 1992. Moreover, Jones has also recently updated the state of modeling of growth and microstructure in rapid solidification. Rapid solidification is defined as the rapid extraction of thermal energy during the transition from the liquid state at high temperatures, to solid material at ambient temperature. The rapid extraction of thermal energy permits large deviations from equilibrium that, offers the advantages of the extension of solid solubility, a reduction in grain size, a reduction in both the number and size of segregated phases, and production of non-equilibrium alloy phases which leads new properties. In practice, there are numerous ways to produce rapidly solidify materials. The processing methodology might involve a moving or stationary substrate or jets of gas or liquid. The main rapid solidification processes are atomization, ribbon or foil casting and spray deposition. -XXXll- In atomization, a fine dispersion of droplets is formed when molten metal is impacted by a high energy fluid (gas or liquid). Atomization occurs as a result of the transfer of kinetic energy from the atomizing fluid to the molten metal. There are three major commercial methods of atomization to produce metal powders : Two-fluid atomization using water or gas, vacuum or soluble gas atomization, and the rotating electrode process. Two-fluid atomization is the predominant technology and accounts for over 95% atomization capacity worldwide. The melt spinning process was developed during the 1960s to produce amorphous ribbons of several alloys. A process that is widely practiced to manufacture continuous rapidly solidified ribbon is chill block melt spinning (CBMS), in which a stream of molten alloy is brought into contact with a rapidly moving substrate surface. A molten alloy puddle form on the moving substrate, resulting from continued melt impingement; this puddle serves as a local reservoir from which ribbon is continuously formed and chilled. The most common substrate surfaces are the inside of drums or wheels, the outside of wheels, and belts. In the planar flow method, the melt ejection crucible is held very close to the moving substrate surface, which causes the melt to be simultaneously in contact with the nozzle and the moving substrate. This entrained melt flow damps perturbations of the melt stream and thereby improves ribbon geometric uniformity. The initial quench rate and final ribbon geometry in free jet melt spinning depends on the nature of the melt puddle. Melt spinning makes possible the production of long narrow ribbons with cooling rates as high as 106oK/s. Generally casting of the ribbon can be divided into three phases. The first stage is the formation of a stationary puddle between nozzle and the surface of the cooling wheel. This phase depends on the size of the nozzle orifice, the nozzle overpressure, the distance between the nozzle exit and the surface of the cooling wheel and on the physical properties of the melt. The melt viscosity and its surface tension are basic parameters influencing the shape of the puddle. Ribbon formation is the second stage and depends on the kinetics of the solidification process. The third stage represents the cooling of the ribbon from the temperature of solidification to room temperature. Spray deposition, Osprey Process, is, a patented, rapid solidification technology for producing semi-finished, spray-formed products in the form of tubes, billets, flats, and compound products in a single integrated operation. A molten metal stream is converted into a spray of droplets by gas atomization; the droplets cool rapidly in flight prior to impacting in a semi-solidified condition a 'collector', where the droplets re-coalesce to form a spray-deposited preform. The shape of preform is determined by movements of both the atomized spray and the collector and the size and shape of the collector. In addition to the production of high alloy ingot metallurgy materials and powder metallurgy type alloys, the Osprey Process is also being used to develop unique materials and products. These include: (a) new alloy compositions, by taking advantage of rapid solidification; (b) Metal matrix composites, by injecting ceramic particulate into the atomized spray during the deposition process; (c) 'in- situ' alloying, by creating reactions with the atomizing gas or the injected particulate; (d) material for semi-solid forming, by creating unique, fine grain- -XXXlll- sized, non-dendritic microstructures; and (e) compound or clad products, by spray- depositing a cladding of a relatively expensive alloy (e.g. stainless steel) an to a lower cost alloy (e.g. mild steel bar) to produce a clad, or compound, bar or billet. Recently, in the transport industries small-sized and low weight, as well as high-performance, machines are being demanded for energy saving, cost reduction and convenience. Stronger, lighter and more economic structural materials are needed for these machines and the aluminum alloys have been paid close attention as potential candidates. Aluminum alloys produced via the conventional ingot metallurgy route are not suitable for use in such ports because of their lack of elastic modulus, coefficient of thermal expansion, strength, heat resistance and wear resistance. The production of Al-Si-Fe-X alloy system by rapid solidification is an important step in this research field. A high amount of Si and large Si particle size are detrimental for extrudability, machinability and mechanical properties. Rapid solidification technologies overcome the defect in Al-Si-Fe-X alloys because finer Si particles are formed by the rapid solidification. Fe and Ni addition in Al- Si-Fe alloys enhance the strength at elevated temperature. 'X' element addition such as Cu, Mg, provides developing of expected properties. Hypereutectic Al-Si-Fe spray-cast alloys exhibit properties similar to those of metal-matrix composite: High young's modulus and low coefficient of thermal expansion. High fatigue properties coupled with high modulus, good high temperature behavior and low thermal expansion, allow their use for applications in the automotive industry. In opposition to MMC's, these materials present the advantage of easy recycling and easy machinability as it is the use for the conventional Al-Si alloys. To contribute to the Turkish automotive industry and research of the rapid solidification and spray deposition effects on the structure properties of Al-Si-Fe-X alloy system, the alloy system which obtained by conventional casting process, produce by some rapid solidification methods, and their properties of solidification, heat treatment, and hardness were investigated. The microstructure, mechanical and heat treatment properties of Al-XSi- 3.3Fe and Al-XSi-3.3Fe-l.7Cu-0.5Mg alloys which produced by conventionally cast, melt-spinning, atomization, and spray deposition process. Al-7.6Si-3.3Fe, Al- 9.8Si-3.3Fe, Al-20Si-3.3Fe, Al-8.4Si-3.3Fe-l.7Cu-0.5Mg, Al-10.4Si-3.3Fe-l.7Cu- 0.5Mg, Al-20Si-3.3Fe-l.7Cu-0.5Mg alloys were prepared using induction melting. Remelting of the as-cast alloys was performed in a graphite crucible affiliated with a H65 type 350kHz high frequency furnace to ensure the homogeneity of compositions. Melt-spun ribbons were produced by impinging a jet of molten alloy onto the cylindrical surface of a polished copper wheel with a diameter of 1 80 mm rotating at 3000rpm. The dimensions of the as-produced ribbons were 6- 10mm wide, 70- 1500mm long and 75- 125 urn thick. During the spray deposition and atomization experiments a cyclindrical stainless steel chamber with a diameter of 400mm. and length of 600mm., was used. To deposit of atomized droplets, a copper disk that was 200mm. diameter and cooled with water continuously, was placed on a hydrolic carrier in the chamber. In the experiments of melt-spinning, atomization and spray deposition, the temperature of the melt before ejection was 100°C above the liquidus temperature. -XXXIV- Optical and scanning electron microscopy observances were performed in an Olympus MG model optical microscope and a JEOL JSM-T330 scanning electron microscope, respectively. The operating voltage of SEM was 20kV. XRD measurements were carried out in a Philips PW3710 X-ray diffractometer using CoKa radiation, in which a PC-APD diffraction software was installed. The samples used for microstructure observances were prepared by mechanical polishing followed by etching in a 0.5% HF solution. The heat treatment of the melt-spun ribbons was performed in a Heraus resistance furnace, in which the temperatures and heating rate were automatically controlled. The thermal analysis was performed using a Perkin-Elmer DSC-7 differential scanning calorimeter with a heating rate of 10°Cs-1. Measurements of microhardness were conducted in a Wolpert microhardness tester. When the silicon content was 7.6wt%., and 9.8 wt%., in Al-XSi-3.3Fe alloys, conventionally as-cast, the microstructure consisted of aluminium matrix, and rod shaped intermetallic phases which contained aluminium, silicon and iron elements. In Al-XSi-3.3Fe-l.7Cu-0.5Mg alloys, conventionally as-cast, when the silicon content was 8.4wt%., and 10.4 wt%., the microstructure consisted of aluminium matrix, and rod shaped intermetallic phases which contained aluminium, silicon, iron, and copper elements. In Al-20Si-3.3Fe and Al-20Si-3.3Fe-l.7Cu- 0.5Mg alloys conventionally as-cast, aluminium matrix, rod shaped intermetallic phases, and also coarse silicon plates which contained approximately 98 wt%., silicon element, were determined. For the melt-spun ribbons with Al-Si-Fe alloys, mean cooling rate was approximately 106-107degC/sec. From the SEM investigation of the melt-spun ribbons with Al-Si-Fe alloys, it was determined that Al-20Si-3.3Fe and A1-20Sİ- 3.3Fe-1.7Cu-0.5Mg alloys which were hyper-eutectic under equilibrium conditions, could have a hypo-eutectic structure as a result of undercooling through rapid-quenching, and these alloys showed a primary aluminium-rich solidification structure. The SEM photographs, given for the wheel and upper sides of the melt- spun ribbon with Al-9.8Si-3.3Fe alloy, it could be seen that the structure at the upper side was coarser than at the wheel side. The SEM investigations of the aged melt-spun ribbons with Al-XSi-3.3Fe-l.7Cu-0.5Mg exhibited the average size of the spherical precipitates increased from 0.4p.m. to 0.94um. with increase in silicon content from 8.4% to 20%. From the stereo microscopic investigations of the melt-spun ribbons, the serrated ribbon edge formation was observed. In all cross-section surface micrograph of the melt-spun ribbon samples, at the wheel-side, discontinuously, a white zone did not respond to etching, was observed. X-ray diffraction spectrum of the melt-spun ribbons with Al-7.6Si-3.3Fe, Al-9.8Si-3.3Fe, Al-8.4Si-3.3Fe-l.7Cu-0.5Mg, as-cast, was similar to that of typical aluminium. Any diffraction line came from an intermetallic phase between aluminium, silicon and iron did not detected. With comparing X-ray investigations of naturally aged melt-spun ribbons and the melt-spun ribbons, as-cast, the peaks indicated Al-matrix phase, silicon phase and P-Al^FeSi phase, were detected from the X-ray spectrum of naturally aged melt-spun ribbons. In DSC investigation for the melt-spun ribbons of Al-XSi-3.3Fe and Al-XSi-3.3Fe-l.7Cu-0.5Mg alloys, an exothermic reaction was observed in range of 310-430°C that was considered to be caused largely by the precipitation of the silicon phase in the aluminium matrix. The endothermic peaks, observed in range of 550-620°C, were considered to mean the melting of aluminium. The mean cooling rate of the powders with Al-Si-Fe alloys for +500um., was 102degC/sec. And the cooling rate of the powders in range of 500um.- lOOum. was estimated to be in the range of 102 to 104degC/sec. It was determined that the particle shape of the powders with Al-Si-Fe alloys for +500um., and -90 um. was similar to spherical and tear-drop respectively. X-ray diffraction spectrum of the powders with Al-7.6Si-3.3Fe, Al-9.8Si-3.3Fe, Al-8.4Si-3.3Fe-l.7Cu-0.5Mg, and Al-10.4Si-3.3Fe-l.7Cu-0.5Mg alloys, as-atomized, was similar to that of typical aluminium. From the X-ray diffraction spectrum of the powders with hyper-eutectic Al-20Si-3.3Fe and Al-20Si-3.3Fe-l.7Cu-0.5Mg alloys, the peaks indicated Al-matrix phase, silicon phase, metastabil intermetallic 8-Al4FeSi2 phase and P-Al5FeSi phase were detected. The mean cooling rate of spray deposited Al-Si-Fe preforms was estimated to be between 103 and 104degC/sec. The optic microscope investigation of the spray-deposited preforms with Al-Si-Fe alloys exhibited that the obtained microstructure was homogeneous. The most of spherical colonies with size of 25- 50um. characterized by the finer and more densely distributed silicon and intermetallic phases in their interior. From the X-ray investigation of the spray- deposited preforms with Al-Si-Fe alloys, it was determined that the peak intensities of silicon phase and P-Al5FeSi phase increased with increase in silicon content. The hardness values of Al-XSi-3.3Fe and Al-XSi-3.3Fe-l.7Cu-0.5Mg alloys which produced by conventionally cast, melt-spinning, atomization and spray deposition process, increased with increase in silicon content. The hardness of the melt-spun Al-XSi-3.3Fe alloys were different from those of the same alloys conventionally cast, and up to approximately X3 enhancements were observed. For the melt-spun Al-XSi-3.3Fe-l.7Cu-0.5Mg alloys, up to approximately X2 enhancements were observed. It was determined that the hardness of the powders with Al-Si-Fe alloys was related to the particle size and silicon content and the hardness ratio of the powders to conventionally cast samples increased with decrease in the particle size.