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 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.

The influence of a low Ti addition (∼0·01%) on the hot ductility of Nb containing HSLA steels has been examined. For conventional cooling conditions in which an average cooling rate from the melting point to the test temperature was used, the ductility decreased markedly with the addition of Ti. However, when cooling conditions after melting were more in accord with the thermal heat treatment undergone by the strand during continuous casting, i.e. cooling is fast to begin with, reaches a minimum and then reheats, after which the temperature falls more slowly to the test temperature, the Ti addition was found to be beneficial.

The influence of B on the hot ductility of high Al, Ti containing twinning induced plasticity (TWIP) steels has been examined. It was established that provided the B was fully protected by adding sufficient Ti to combine with all the N, then B could segregate to the austenite grain boundaries and improve ductility. This improvement was particularly marked for the temperature range of 700–900°C, the range in which the straightening operation often takes place in continuous casting. Of most importance in the present work has been the detection of B at the boundaries using a secondary ion mass spectrometry technique. The cooling rate from the reheating temperature of 1250°C to the tensile testing temperature range of 700–1200°C was 60 K min−1, but it is likely that slower cooling rates ≤25 K min−1, more in keeping with the secondary cooling rate on continuous casting, will give even better ductility. Ti additions in themselves are beneficial to the hot ductility of these steels as precipitation of AlN at the austenite boundaries is avoided, but only if the cooling rate is sufficiently slow to allow the TiN particles to coarsen. However, to ensure freedom from cracking, an addition of B is also required.

A variety of heating and cooling programmes have been examined for plain C-Mn and high strength low alloy steels to examine their suitability in a hot tensile test for assessing the likelihood of transverse cracking occurring in the straightening operation. A tensile test temperature of 800uC was chosen for comparison, this being the temperature that generally results in poor ductility. For steels with 1?4-1?75%Mn, the simple procedure of heating to y1300uC to take all the microalloying additions into solution, followed by cooling to the test temperature, was found to be the easiest and most suitable. For high strength low alloy steels containing Ti and low Mn, high S steels, melting is required. For these steels, it is also advisable to have both primary rapid cooling followed by a slower secondary cooling stage, simulating more accurately the actual industrial operation. The addition of thermal oscillations to simulate slab roll contact makes the cycle even closer to the commercial process and generally led to a small decrease in reduction in area values. Because of its complexity, this latter method would not be generally recommended for steels showing wide trough behaviour (high Mn, peritectic C steels), and melting followed by primary and secondary cooling is sufficient. For narrow troughs (low Mn, low C steels), which require melting and where the minimum ductility will be at a temperature of .800uC, the more complex procedure will be required. It will be necessary to obtain the full hot ductility curve using a cycle that, as well as melting and having primary and secondary cooling, also incorporates commercial thermal oscillations or at least some limited thermal oscillations.