A novel ultra-high strength precipitation hardened martensitic steel with balanced ductility and creep resistance has been developed. It utilises a unique combination of nanometre scale intermetallic precipitates of Laves phases and β-NiAl to achieve such properties. The mechanical properties of this steel were assessed by tensile and creep testing. With different heat treatments, this steel showed a remarkable combination of mechanical properties: yield strength of >1800 MPa, ultimate tensile strength of ~ 2000 MPa, tensile ductility up to ~8% at room temperature and creep rupture life > 2,000 hours under 700 MPa stress at 500 °C. The microstructures at different length scales were characterised using scanning / transmission electron microscopy and atom probe tomography. The austenisation and ageing temperatures were found be the key factors determining the microstructural development and resulting mechanical properties. Large primary Laves phase precipitates formed at lower austenisation temperatures resulted in reduced creep strength; whilst the small difference (20 °C) in ageing temperatures had significant impact on the spatial distribution characteristics of β-NiAl precipitates. Lower ageing temperature produced much smaller but more uniformly distributed β-NiAl precipitates which contributed to the higher observed yield strength. It is clear from this study that whilst this novel alloy system showed great potentials, careful design of heat treatment is still required to achieve balanced mechanical properties to meet the service requirements in aerospace propulsion systems.

In the current study, dislocation activity and storage during creep deformation in a nickel based superalloy (Waspaloy) was investigated, focussing on the storage of geometrically necessary (GND) and statistically stored (SSD) dislocations. Two methods of GND density calculations were used, namely; EBSD Hough Transformation and HR-EBSD Cross Correlation based methods. The storage of dislocations, including SSDs, was investigated by means of TEM imaging. Here, the concept of GND accumulation in soft and hard grains and the effect of neighbouring grain orientation on total dislocation density was examined. Furthermore, the influence of applied stress (below and above Waspaloy yield stress) during creep on deformation micro-mechanism and dislocation density was studied. It was demonstrated that soft grains provided pure shear conditions at least on two octahedral (111) slips for easy dislocation movement reaching the grain boundary without significant geometrically necessary accumulation in the centre of the grain. Hence, the majority of the soft grains appeared to have minimum GND density in the centre of the grain with high GND accumulation in the vicinity of the grain boundaries. However, the values and width of accumulated GND depended on the surrounding grain orientations. Furthermore, it was shown that the hard grains were not favourably oriented for octahedral slip system activation leading to a grain rotation in order to activate any of the available slip systems. Eventually, (i) the hard grain resistance to deformation and (ii) neighbouring grain resistance for the hard grain reorientation caused high GND density on a number of octahedral (111) slip systems. The results also showed that during creep below the yield stress of Waspaloy (500 MPa/700C), the GND accumulation was relatively low due to insufficient microscopic stress level. However, the regions near grain boundaries showed high GND density.Whereas, in addition to the movement of pre-existing dislocations (SSD and GND) at higher mobility rate under 800 MPa/700C above yield creep condition, large numbers of dislocations were generated and moved toward the grain boundaries. This resulted in much higher GND density but narrower width of high intensity GND near the grain boundaries. It is concluded that although GND measurement by means of EBSD can provide a great insight of dislocation accumulation and its behaviour, it is critical however to consider SSD type which is also contributes to the strain hardening of the materials.

Moat, In situ observation of strain and phase transformation in plastically deformed 301 austenitic stainless steel, (2016), AbstractTo inform the design of superior transformation-induced plasticity (TRIP) steels, it is important to understand what happens at the microstructural length scales. In this study, strain-induced martensitic transformation is studied by in situ digital image correlation (DIC) in a scanning electron microscope. Digital image correlation at submicron length scales enables mapping of transformation strains with high confidence. These are correlated with electron backscatter diffraction (EBSD) prior to and post the deformation process to get a comprehensive understanding of the strain-induced transformation mechanism. The results are compared with mathematical models for enhanced prediction of strain-induced martensitic phase transformation.

Diffraction peak profile analysis (DPPA) is a valuable method to understand the microstructure and defects present in a crystalline material. Peak broadening anisotropy, where broadening of a diffraction peak doesn't change smoothly with 2θ or d-spacing, is an important aspect of these methods. There are numerous approaches to take to deal with this anisotropy in metal alloys, which can be used to gain information about the dislocation types present in a sample and the amount of planar faults. However, there are problems in determining which method to use and the potential errors that can result. This is particularly the case for hexagonal close packed (HCP) alloys. There is though a distinct advantage of broadening anisotropy in that it provides a unique and potentially valuable way to develop crystal plasticity and work-hardening models. In this work we use several practical examples of the use of DPPA to highlight the issues of broadening anisotropy. of 31where, D is the crystal size, K Sch is the Scherrer constant (often taken as 0.9 for spherical crystals), g is the reciprocal of the d-spacing of a peak, f m is a function related to the arrangement of dislocations (and represents how the arrangement of groups of dislocations influence their total strain), ρ is the dislocation density and C hkl is the average contrast factor of dislocations in grains that contribute to the hkl diffraction peak. In Equation 1 and other DPPA approaches, C hkl is assumed to be the only term that accounts for broadening anisotropy and its value is required to be able to obtain the dislocation density.TEM to be able to identify details of dislocations [15,16], especially in cases when g.b=0 does not lead to vanishing contrast or due to practicalities of rotating a sample. In the same manner, the diffraction broadening caused by a dislocation varies depending on the diffraction vector, and can be calculated by modelling the dislocation's displacement field. The contribution of a dislocation's displacement field to the broadening of different diffraction profiles is through what is known as the contrast (or orientation) factor of dislocations. The contrast factor of an individual dislocation is dependent on the angles between the diffraction vector and the vectors that define the dislocation: it's Burgers vector (b), slip plane normal (n), and dislocation line (s). The contrast factor of an individual dislocation can be calculated using the computer program ANIZC [17]. For example, the edge dislocation in Figure 2 has a contrast factor of 0.46 when g is [110] and parallel to the dislocations Burgers vector. The value is 0.00 when g is parallel to the slip line ([11 2]) and 0.05 when g is parallel to the slip plane normal ([1 11]).

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