The hot ductility of twin induced plasticity (TWIP), 0·6 wt-C steels containing 18–22 wt-Mn with N levels in the range 0·005–0·023 wt- and the Al additions either low (<0·05 wt-) or high (1·5 wt-) has been examined. Little change in ductility occurred in the temperature range 1100–650°C as the structure was always fully austenitic. Ductility was generally poor (<40), reduction of area values, the best ductility at the higher Al level being given by the steel with the lowest N and S levels. Because the steel is fully austenitic, the ductility is solely dependent on that for unrecrystallised austenite. Therefore, to avoid transverse cracking the volume of second phase particles should be kept to a minimum, i.e. the N should be low to reduce the amount of AlN that can be precipitated out and the S level should be as low as possible to limit the amount of MnS inclusions. Metallographic and TEM studies were carried out and the poor ductility was found to be due to extensive precipitation of AlN at the austenite grain boundaries. Increasing the cooling rate from the melting point to the test temperature from 60 to 180°C min−1 or introducing an undercooling step both led to even worse ductility.

The hot ductility of a high Al, twin induced plasticity austenitic steel with S levels of 0·003, 0·010 and 0·023 has been examined after heating to 1250°C and cooling at 60 K min−1 to test temperatures in the range of 700–1100°C. Ductility as measured by reduction of area ( R of A) was very poor, (∼20 R of A) in the two higher S steels in the temperature range of 950–1100°C but was better at lower temperatures of 700–900°C reaching 30–40 R of A. For the very low S steel, ductility was similar for the temperature range of 700–900°C but improved in the higher temperature range of 950–1100°C to 50–55 R of A. In the higher S steels (0·01 and 0·023S), ductility was poor because of the presence of long, coarse dendritic AlN rods situated at both the dendrite, but, more importantly, the austenite grain boundaries, the latter being particularly detrimental for encouraging intergranular failure. AlN seems to need MnS inclusions as nuclei in the melt for it to precipitate in the form of these detrimental long rods, and a reasonable volume fraction of MnS (equivalent to having more than 0·003S present) is required for this to happen. If the S level is low (⩽0·003), there are only a few MnS nucleation sites, and the AlN precipitates out in the form of coarse hexagonal plates. These precipitates end up mainly within the austenite matrix and have little influence on the hot ductility. It is therefore very important commercially in these high Al steels to limit the S level to ⩽0·003 to avoid transverse cracking.

The hot ductility of Nb/V containing high Al, twin induced plasticity (TWIP) steels has been examined over the temperature range 650–1150°C after melting and after ‘solution treatment’. Previous work had shown that the hot ductility is poor for the 1·5 mass-Al, TWIP steel due to precipitation of AlN at the austenite grain boundaries, the depth of the trough being similar to that for an X65 grade pipeline steel but with the trough covering a much wider temperature range. Adding Nb and V made the ductility even worse due to the additional precipitation of NbCN and VN. Very low reduction of area values, 10–20 were obtained in the temperature range 700–900°C. Increasing the cooling rate to the test temperature resulted in even worse ductility. The ductility of these steels after ‘solution treatment’ is similar to that obtained after melting but when the cast was hot rolled followed by ‘solution treatment’ and cooling to the test temperature ductility improved due to grain refinement.

High Al, twinning induced plasticity (TWIP) steels have high tensile strengths and excellent ductility but low yield strengths compared to other advanced high strength steels and this has limited their application. A.Vanadium addition is a possible answer but cracking on casting is a worry. Hot tensile tests were therefore performed on 1.5% Al, TWIP steels with vanadium levels approximately 0.05–0.7% (all wt-%). Only the 0.05% vanadium steel gave acceptable hot ductility but the room temperature yield strength was too low. In contrast, the Ti–B high 0.5% V steel which was as well as able to give a high yield strength due to precipitation hardening by vanadium carbide gave better ductility by boron segregating to the boundaries and strengthening them.

The addition of ∼0·002B and ∼0·04Ti as microalloying additions to improve the poor hot ductility and high risk of cracking on continuous casting of high Al containing twinning induced plasticity (TWIP) steels has been examined. Tensile specimens were either cast in situ or heated to 1250°C before cooling at 60 K min−1 to test temperatures in the range 700–1100°C and strained to failure at 3×10−3 s−1. For tensile specimens reheated to 1250°C, the presence of B with sufficient Ti to combine with all the N improved ductility over the temperature range of 700–950°C, the reduction in area (RA) values being >40. For the higher strength more complex high Al, TWIP steels having Nb present, there was no improvement in ductility with a similar B and Ti addition, when the average cooling rate after melting to the test temperature was 60 K min−1. Reducing the cooling rate to 12 K min−1 resulted in the RA values being close to the minimum required to avoid transverse cracking throughout the temperature range 800–1000°C. Using these additions of B and Ti, transverse cracking was found not to be a problem when continuously casting these high Al containing TWIP steels.