Yield (engineering) | EPFL Graph Search

In materials science and engineering, the yield point is the point on a stress-strain curve that indicates the limit of elastic behavior and the beginning of plastic behavior. Below the yield point, a material will deform elastically and will return to its original shape when the applied stress is removed. Once the yield point is passed, some fraction of the deformation will be permanent and non-reversible and is known as plastic deformation. The yield strength or yield stress is a material property and is the stress corresponding to the yield point at which the material begins to deform plastically. The yield strength is often used to determine the maximum allowable load in a mechanical component, since it represents the upper limit to forces that can be applied without producing permanent deformation. In some materials, such as aluminium, there is a gradual onset of non-linear behavior, and no precise yield point. In such a case, the offset yield point (or proof stress) is taken a

Microstructure of single crystal Ni-based superalloys during blade manufacturing process and creep deformation

Stéphane Pierret

The aim of the thesis is to obtain more understanding about the influence of plastic deformation on the microstructural changes observed in single crystal (SX) Ni-based superalloys using diffraction techniques. This work was organised around two main problematics: (1) rafting in turbine blades after the manufacturing process and (2) rafting and mosaicity evolution in two SX alloys with different composition during creep deformation. SX Ni-based superalloys are first choice materials for turbine blade application in land-based gas turbines and aero-engines. Superior mechanical properties of these materials at high temperatures are achieved due to the optimisation of the alloying composition and the microstructure. SX Ni-based superalloy turbine blades solidify by dendritic solidification with dendrites aligned along the directions of lowest Young modulus (interdendritic spacing around 300μm). At the end of the manufacturing process, the microstructure consists of high volume fraction of cubic γ’ precipitates (around 500nm) coherently embedded in the γ matrix. However specific parts of engine-ready turbine blades exhibit directionally coarsened γ’ precipitates in a preferred direction (rafting) without the influence of an applied load. It is of prime importance to understand the origin of rafting prior to the introduction of the blades in the engine since rafting can alter the mechanical properties of the alloy. Chemical segregations and the introduction of plastic deformation prior to thermal treatments were identified as potential driving force for rafting during load-free high temperature annealing. Most of the SX Ni-based superalloys used in the current blade production in Alstom contain significant amount of refractory elements which are expensive and increase the density of the alloy. To answer the demand of cheaper and lighter components, Alstom launched a development program to design novel SX superalloys. One of these projects led to the production of a Re-free superalloy, so-called MD2. Extensive mechanical investigations on MD2 are on-going and compared to the behaviour of the MK4HC alloy (3wt% Re) currently used in turbine blade production. One of the major differences in terms of mechanical properties between Re-containing and Re-free SX alloys is the faster kinetics of coarsening and rafting during creep deformation at high temperature in the latter alloys. It prevented so far their application in the hottest areas in gas turbines. It is therefore of utmost importance to acquire more knowledge about the mechanisms associated with the microstructure evolution during creep of the new MD2 SX alloy in comparison with the MK4HC alloy. The evolution of the lattice parameter misfit in different directions with respect to the applied load during creep deformation was associated with the evolution of the microstructure from cubic to elongated γ’ precipitates It could be shown that rafting occurs after the cooling stage following the solution heat treatment (SHT) during the manufacturing of the blade with investigations of the microstructure. The influence of chemical segregations on rafting was excluded with the investigation of the blade alloy composition in different parts of the blade cooled from the SHT and diffraction measurements in stress-relieved combs prepared from the same blade. The formation of a rafted microstructure during the manufacturing process was associated with the formation of plastic strains during cooling after the SHT. This was confirmed after gathering results from investigations of the microstructure, simulation of the cooling stage and neutron diffraction measurements on the blade. The investigations of the microstructure performed on the engine-ready blade revealed an inversion of the γ’ raft orientation below (parallel to [001]) and above the platform of the blade (perpendicular to [001]). This was attributed to an inversion of the sign of residual strain between the previously mentioned positions as revealed by neutron diffraction measurements performed along the airfoil of the blade cooled after the SHT. Microscopy also showed that the rafted microstructure was more pronounced below the platform (long γ’ rafts) than above the platform (shorter γ’ rafts with retained cubic precipitates). The concept of the threshold plastic strain necessary to induce rafting during annealing subsequent to plastic deformation was used to describe the difference in rafted microstructure. It is suggested that plastic strains largely exceed the threshold below the platform leading to fully formed rafted microstructure. On the other side, plastic strains of the order to the threshold strain are suggested above the platform in order to explain the partially formed rafted microstructure Neutron diffraction measurements revealed that the lattice parameter misfit was larger in the MK4HC than in the MD2 alloy at 20°C in the unloaded state. The larger misfit in the MK4HC alloy was attributed to the larger concentration of elements dissolving in the γ phase and to the presence of Re. During creep deformation, coherency stresses are relieved at the γ/γ’ interface in the γ channels perpendicular to the applied load by dislocations. On the other hand, dislocations increase the internal stress at the γ/γ’ interface in the γ channels parallel to the applied load. This leads to the rafting of the γ’ precipitates perpendicular to the applied load. This was observed in the MK4HC. The absence of clear misfit evolution in the MD2 alloy was attributed to the anomalous microstructure at the beginning of the creep deformation. The mosaicity of the MK4HC and MD2 alloys revealed crystallographic misorientation between neighbouring dendrites up to 0.5°. The misorientation was associated with the presence of dislocations in the regions between the dendrites formed during the dendritic solidification of the alloys from existing work. It was shown that the composition of the SX alloys and short periods at high temperature in the unloaded state had only little influence on the misorientation between neighbouring dendrites. Upon creep loading to 180MPa, the mosaicity markedly changed in the two SX alloys independent of the creep temperature. In particular, the number of single misoriented domains measured by X-ray diffraction decreased. This was attributed to the fast propagation of the dislocations present in the interdendritic structure inducing a rotation of the individual dendrites. During creep deformation, the mosaicity increased as visible from the broadening of the maximum spread in misorientation. The formation of inhomogeneous plastic strain fields in the γ/γ’ microstructure is believed to create new mosaic blocks in the dendritic structure of the alloys leading to the observed broadening of the mosaicity. The influence of plastic deformation on the evolution of the mosaicity was demonstrated. The mosaicity significantly changed during creep of the MD2 at 1050°C/180MPa and 1000°C/180MPa whereas no significant evolution could be measured during creep at 950°C/180MPa. This was attributed to a temperature dependence of the yield strength of the MD2 alloy. It therefore induces difference in dislocation propagation at 950°C compared to 1050°C and 1000°C which retards the broadening of the mosaicity during creep at 950°C. Furthermore, the mosaicity formed during creep deformation was retained upon unloading and cooling of the SX samples to room temperature.

EPFL2012

In-situ Laue Diffraction During Compression of Directionally Solidified Mo Micropillars

Julien René André Zimmermann

With recent developments in micro and nano-technologies, mechanical components such as those used in medicine or electronics tend to be miniaturized, requiring a new set of testing techniques to study their reliability and performance. One of the new methods is micro-pillar compression testing, which allows the study of mechanical properties of confined volumes precisely shaped by focused ion beam (FIB) milling from sometimes complex microstructures. This technique has garnered a great deal of interest in recent years; on one hand because of its apparent simplicity, on the other hand because it revealed an unexpected size dependence in the flow stress of single crystals when the pillar diameter was reduced below a few tenths of a micrometre. The flow stress of FIB milled face centred cubic (fcc) micro-pillars appears to follow a power law as a function of their size. In contrast, non-FIB milled, body centred cubic (bcc) eutectic Molybdenum (Mo) alloy pillars prepared by directional growth exhibit, in their pristine state, whisker-like behaviour without a size effect. Furthermore, after FIB milling or moderate pre-deformation, the whisker character is lost and the typical scatter of the flow stress usually observed for FIB milled pillars is reproduced. After large pre-deformation the scatter vanishes and the pillars exhibit no size effect. These findings show that 1) the deterministic parameter in the size effect can be correlated with the initial microstructure of the pillar, and 2) FIB milling influences this microstructure. This study aims to link the deformation modes occurring in small-scale pillars with their initial microstructure by studying directionally solidified eutectic Mo-alloy pillars with a diameter of 1 micrometre. In-situ Laue diffraction experiments during the compression of such defect-free pillars show that with readily-available dislocation sources from a concave pillar top, the {112}[111] slip system with the highest Schmid factor is initially activated and not the allegedly assumed {110}[111] slip system that is typical for bcc metals. Post-mortem slip trace analysis and TEM observations confirm this finding. Moreover, 3D atom probe measurements and TEM images show that this behaviour can be linked to the presence of both alloying elements and fcc nano-precipitates found in the bcc Mo-alloy. Both can lead to a decrease of the critical temperature and the observation of an fcc-like deformation. Additionally, it is shown that both FIB milling and pre-deformation introduce defects in the initially dislocation-free pillars. In addition to this, pillars prepared under similar conditions yield quite different microstructures. The angle between the focused ion beam and the crystal orientation is found to play an important role in the creation of defects. It can also be shown that pillars with high initial dislocation density deform bulk-like and exhibit strain hardening, while pillars with low initial dislocation content demonstrate a large scatter in the flow and yield stress and exhibit large strain burst at later deformation stages. Furthermore several slip systems are activated before the flow stress is reached for pillars containing initially low dislocation content. Interestingly the first activated slip system can be linked to the initial streaking direction; slip is observed on the {112} plane that contains excess dislocations. For both the as-prepared pillars and the pillars with low initial dislocation content, the yield seems to be source limited, which is not the case for pillars with high initial dislocation content.

EPFL2011