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Phase Transformation In Metals And Alloys Porter Pdf Free _BEST_

  • Entirely new and extensive treatment of diffusionless martensitic transformations covering athermal and thermally activated martensite in ferrous systems as well as shape memory, superelasticity and rubber-like behavior in ordered nonferrous alloys

phase transformation in metals and alloys porter pdf free

David Porter studied materials science in Cambridge obtaining his Ph.D. in 1976. He subsequently moved to the University of Luleå, where he applied electron microscopy to materials research and taught courses on electron microscopy and phase transformations. He subsequently worked in the research and development departments of the Norwegian aluminum company Årdal og Sunndal Verk and the steel producers Rautaruukki in Finland, and Fundia Special Bar in Sweden. In 2011 he returned to academia as professor of physical metallurgy at the University of Oulu where he has been professor emeritus since 2019.

The results should be regarded as an initial prediction of the phases and transition temperatures. This is due to the fact that the Ta-Cu binary, the Ta-Nb-Cu ternary and the Ti-Ta-Cu ternary systems have not been thermodynamically assessed and thus are lacking in the SSOL5 database. Nevertheless, the predictions given by the calculations are useful as a starting point for alloy development and to guide the experimental work. The Ti-1.7 wt.% Nb-10.1 wt.% Ta-1.6 wt.% Zr (TNTZ) has been modelled previously [11] and gave only α and β phases. The equilibrium phases as a function of temperature were modelled for this alloy with increasing Cu additions (0 wt.% Cu, 1 wt.% Cu, 3 wt.% Cu, 5 wt.% Cu and 10 wt.% Cu). In Figure 1a,b the phase fractions in the alloys with 1 wt.% Cu and 5 wt.% Cu addition are shown. Given the prerequisites mentioned, the Ti2Cu forms at 656 C in the 1 wt.% Cu alloy (Figure 1a) in thermodynamic equilibrium; However, since the mole fraction is very low it is not likely to nucleate due to kinetic reasons. Nevertheless, when the phase fraction of Ti2Cu increases with Cu addition, already at 3 wt.% Cu and here at 5 wt.% Cu (Figure 1b), the phase fraction is considerable. The modelling resulted in the phase transition temperatures given in Table 1, where transus in this case, is the temperature above which the phase is no longer stable.

Mole fraction of phases as a function of temperature for (a) the Ti-Nb-Ta-Zr-1 wt.% Cu alloy and (b) the Ti-Nb-Ta-Zr-5 wt.% Cu alloy. Composition of the alloys can be found in Table 3.

The thermodynamic prediction of phases for the alloys is given in Figure 1a,b as a function of temperature, and predicted HCP-Ti (α) and BCC-Ti (β), and additionally Ti2Cu for Cu additions of 1% and higher. At the heat treatment temperature of 747 C, the calculated mol% of the phases is given in Table 5, where the 10 wt.% Cu alloy was predicted to have no α-Ti phase present, while alloys below 5 wt.% Cu were predicted to have no Ti2Cu phase present.

The 3 wt.% Cu alloy had a microstructure similar to the lower Cu content alloys with thin lathes, but by using backscattered electron imaging, smaller precipitates were discovered at the GBs of the larger α-Ti grains (Figure 3e). These areas were studied further by preparation of focused ion beam (FIB) lamella, STEM-EDS and transmission Kikuchi diffraction (TKD). Regions of Cu-rich precipitates were observed, with adjacent crystals containing Ti and Ta (Figure 4a). The grains with the brightest contrast, which probably was β-Ti considering the heat treatment temperature, were also slightly coarser grained in the 3 wt.% Cu alloy compared to those with lower Cu content. The phases were assigned as a matrix phase of α-Ti, Ti2Cu and a bright phase, where the bright phase could not be assigned to a known crystal phase using TKD (Figure 5). The 5 wt.% Cu alloy had a microstructure of irregular lathes compared to those with lower Cu content (Figure 3). The lathes that formed were not straight-line structures as in the 3 wt.% Cu, but instead lathes disrupted by Cu-rich globules, formed along the length of the bright β-phase.

Comparison of the predicted β-transus temperatures to the experimentally observed values, revealed discrepancies in the data at 829 C, 751 C, 746 C and 744 C for the 1, 3, 5 and 10 wt.% Cu alloys, respectively. The measurements were in good agreement with the calculated values of 746 C and 753 C, for the 10 wt.% Cu and 5 wt.% Cu alloys. However, the discrepancy between the calculated and measured values increased as the Cu content decreased. The reason for the discrepancies could be the absence of the Ti-Ta-Cu system in the database. An additional cause for variance in the discrepancies could be due to the reduction in the effective Cu content, since Cu is bonded in the intermetallic (Ti2Cu) phase, which was identified by diffraction for the 5 and 10 wt.% Cu alloys. A further reason could be that the β-stabilizers of Ta and Nb are soluble in the intermetallic phases, in addition to Cu. It is also uncertain whether the metastable Ti3Cu [24] is present for the lower Cu compositions, thus further research is required.

The 0 wt.% Cu TNTZ was heat treated at 747 C, and predicted to have a microstructure consisting of α (76.4%) and β (23.6%), with a hardness of 135 3 Hv. In a previous study [11] the alloy was found to be a α (50%) and β (50%) alloy with hardness of 340 HVN. The differences found were probably due to various forging treatments of the alloy in the previous study [11]. The addition of 1 wt.% Cu did not cause a third phase to precipitate, presumably due to the fact that the Ti2Cu phase only forms at temperatures lower than 747 C (Figure 1a). Therefore the 0 wt.% Cu and 1 wt.% Cu alloys are confirmed as two-phased materials via diffraction (Figure 2) and microscopy studies (Figure 3).

The applications of the chromium ferritic stainless steel AISI 410S have been considerably increased in the last years in many technical fields as chemical industries and oil or gas transportation. However, the phase transformation temperatures are, currently, unknown for this alloy. The aim of this work is to determine the alpha to gamma transformation temperatures of the AISI 410S alloy in different cooling conditions and to analyze them using continuous cooling theory. In order to achieve different cooling rates and thermal conditions, two complementary techniques were used: Bridgman furnace crystal growth and laser remelting technique. The measured solidification temperature was around 1730 and 1750 K. Plate-like and dendritic austenite precipitates were obtained in solid-state phase using growth rates between 5 and 10 µm/s in directional growth experiments. Only plate-like austenite phase was observed in the experiments using growth rates above 100 µm/s. The appearance of dendrites, with the consequent segregation of the elements, can be previously determined by the microstructure modeling currently proposed. Massive austenite can be produced from 0.3 to 10 mm/s rates at temperatures between 1100-1300 K. The structure might be less sensitive to corrosion because this phase is produced without microsegregation.

The applications of chromium ferritic stainless steels have been considerably increased in the last years in many technical fields as chemical industries and oil or gas transportation. Thanks to the combination of its high-corrosion resistance and good mechanical properties1, these alloys can be found in different environments as cargo ships and external architectural facades. Currently, the industries reduce the use of strategic and costly metals, such as Ni, and prefer metals which maintain the corrosion and mechanical properties near to that of austenitic grade stainless steels. For this, the AISI 410S (European grade 1.4003) with 10.5-12.5% Cr and less than 1% Ni was developed2,3.

Although, the final microstructure is known to be composed by variable amounts of ferrite and martensite, still the 410S duplex microstructure is not well understood. Especially, the effect of cooling rate on the formation of the austenite phase and the microsegregation pattern linked to the ferrite-austenite transformation are even unclear4. It is known that AISI new requirements of the petroleum refining industry5-8 need further research about the microstructure evolution during solidification and solid-state transformations of these steels.

Most of the phase transformation studies in ferritic/austenitic steels9-111, were based on isothermal treatments. This method usually requires a large number of specimens for a complete description of reactions. On the other hand, crystal growth techniques like Bridgman12 can offer a very controlled way to verify the influence of processing conditions on the kinetics of phase transformations in a single sample. In this case, the growth is imposed by the continuous displacement of isotherms (isovelocity) when the sample is displaced on the vertical axis from the equipment furnace to a quenching medium. Thus, the transformation interface is constrained to assume a given growth morphology and temperature. Phase growth studies using directional solidification have been carried out by a number of authors, like Trivedi et al. for the Al-Cu system13, Lima and Kurz14 for the Fe-Cr and Fe-Ni systems and Jacot et al.15 for Fe-Co alloys.

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