PHASE TRANSFORMATION IN INTERSTITIAL-FREE, AND LOW DENSITY STEELS

Abstract

Interstitial-free (IF) steel is commercial grade steel for deep drawing applications over the last few decades since in 1980s. The term ‘interstitial free’ refers to the fact that interstitial components are present in a trace quantity (30-40 ppm) to strain the iron lattice, resulting in a softer steel. At present the steel is used regularly for the outer shells of cars, beverage cans, refrigerator enclosure, enamel wares, house hold appliances etc. The primary reason is the extremely high formability, capable of providing complex shape geometry. This steel is reported to have Lankford parameter ≥1.8 and the strain hardening exponent ≥ 0.22, as ideal candidate for automotives. Auto components with immense intricacy, e.g. fenders, bonnet hood or other outer panels today find no alternatives to the IF steels. The interstitial elements like carbon and nitrogen (residual amount) are further scavenged by titanium and niobium in the steel. It imparts the absence of yield point elongation or Lüders bands formation during plastic deformation for a better surface finish. The substitutional elements, in general are also very low (⁓0.7 wt.%). The alloy chemistry thus makes the steel softer than even low carbon steel and less resistant to denting. From time to time, various researchers have intended to address this issue via grain refining, bake hardening, solid solution strengthening, precipitation hardening etc. Interestingly, none of them turns out to be much effective. On the contrary, less attention has been paid so far on the microstructure modifications by phase transformations, which is a key research gap to be addressed in the present study. To execute the work, a commercial grade of IF steel is obtained from Tata Steel, Jamshedpur, India with the commercial-purity of about 99.3% on a weight basis. Thermodynamic prediction of the phase stability has been carried out initially with equilibrium computation of phase diagram by Thermo-calc software. Later the calculation has been verified with dilatometry, and differential thermal analysis (DTA) to ascertain austenitizing temperature for ferrite to austenite formation at around 921 °C. Solid cylindrical specimens (10 mm diameter × 15 mm height) have been machined out from the as-received block in order to carry out heat treatment in a Gleeble 3800® thermo-mechanical simulator. The samples have been soaked at 925-930 °C for 3-5 minutes, followed by different rates of cooling in annealed, normalized and water quenched conditions. Microstructural characterizations have been done under optical microscope, a field emission gun scanning electron microscope with the attachment of an electron backscattered diffraction (EBSD) and high resolution transmission electron microscopy. The TSLOIM software has been used to interpret the EBSD scanned data. The macro-hardness of the investigated samples is measured in the Vickers hardness testing machine. A couple of rectangular blocks (15 mm× 20 mm× 100 mm) have been cut from the as-received sample and undergone same heat treatment regimes in a laboratory scale to fabricate sub-size tensile specimens in accordance with vi the ASTM E8 standard. The fractrographs of the tensile deformed samples have been compared side by side with the EBSD data for the microstructure property correlation. The preliminary investigation suggests that the as-received sample is ferritic (α) at room temperature with a bimodular grain structure varying from 25 ± 5 μm to 240 ± 25 μm for smaller and larger ones. After annealing, nucleation of new grains happens. In normalized (~26 ºC/s) condition bainitic trace can be seen. The water quenched (~ 730 ºC/s) sample exhibits lath type morphology to impart strengthening. With the HRTEM bright field micrographs, fine lath type of features with highly dislocated structures have been observed. In the SAD pattern, evidence of BCC product lattice has also been indexed. The IQ map extracted from EBSD has encountered a highly irregular grain shaped region with ragged boundaries; which in turn imparts a significant difference of image quality values compared to the lath regions in confirming massive ferrite as the duplex phase in accordance with the TTT and CCT. The higher amount of the degree of misorientation, as well as the two-fold increment of dislocation density in the lath region, distinguishes it clearly from the massive ferrite followed by tint etching. The calculated variants on the basis of K-S orientation relationships between parent and product phase as projected in the pole figure from the prior austenite grain designate that the transformation of differently oriented martensite can originate from a single parent austenite grain. However, the maximum achievable tensile strength (UTS) is limited to 447 MPa in the outcome. A further enhancement of strength appears to be difficult by strain induced precipitates around martensite during thermo-mechanical processing, as also a part of the current study. The tetragonality in this steel is essentially ruled out due to lack of carbon (30 ppm) in the solid solution. The role of the hierarchical structure of martensite such as packet, block and sub-block boundaries has been investigated in depth with the EBSD to resolve the controversy on their role to strengthening. The analysis of the data implicates that block, sub-block or martensite variants are less decisive, other than morphologies towards the academic interest. Hereby, translating IF steel into high-strength high-ductile automotive grade steel is apriority in the present work. As the additional part of the work, some aspects of low density steel alloyed with aluminum have been discussed within the purview of the current phase transformation study. It is known that aluminum is the key ingredient in low density steels. It leads to the formation of δ-ferrite in a large fraction. In order to study the influence, two different experimental approaches have been devised. In a first of its kind, the sample has been arc melted in an argon atmosphere. Later the same alloy composition has been ball milled and sintered at 1000 °C by avoiding liquid to δ-ferrite in the system. The comparative analysis lead by Dictra simulation indicates that kappa (κ) carbide is difficult to form when δ-ferrite is prior available in the present alloy system during solidification-casting in particular.