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

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.

On the relationship between the microstructure and the mechanical properties of an ODS ferritic/martensitic steel

Amuthan Ramar

In the quest of materials for the first wall of the future fusion reactor, it has been shown that oxide dispersion strengthened (ODS) ferritic / martensitic (F/M) steels appear to be promising candidates. The inherent good mechanical properties supported by a good thermal conductivity, swelling resistance and low radiation damage accumulation of the base material, such as EUROFER97, are further enhanced by the presence of a fine dispersion of nanometric oxide particles. They would allow in principle for a higher operating temperature of the fusion reactor, above 550°C, which improves its thermal efficiency. In effect, their strength remains higher than the base material at high temperatures. There are however limitations to their application in the reactor, which relate to the intimate link between the microstructure and the mechanical properties. These limitations, such as an increased ductile to brittle transition temperature above room temperature relative to the base material, come directly from the role played by the oxide particles but also from their indirect impact on the overall microstructure, stemming from the procedure of elaboration of these materials. In addition, there is a lack of knowledge on their response to irradiation, which, even though it appears to be promising with a reduced irradiation induced hardening, is not very well understood. In this work we have developed and evaluated ferritic/martensitic steel, EUROFER97, strengthened by a dispersion of yttria particles. (1) The powder metallurgical process used to obtain it was optimized, (2) the basic mechanisms of hardening were identified and (3) the impact of irradiation on its microstructure and mechanical properties were investigated. In order to understand the behaviour of this class of material other oxide dispersion strengthened (ODS) ferritic/martensitic steels based on EUROFER97 and model alloys were also investigated. For the elaboration of the ODS steel the powders of atomized EUROFER97 and yttria are ball milled in a planetary ball mill and the material is compacted by hot isostatic pressing (HIP). The mechanical properties are evaluated by uniaxial tensile test, Charpy test and micro-indentation while the microstructure is investigated by optical microscopy, scanning electron microscopy and transmission electron microscopy. The focused ion beam (FIB) technique is also used. In the study of the ball milling it appears that the optimal time is 20 hours, as the X-ray peak broadening, crystallite size and hardness of the particles remain unchanged at longer times. Yttria nanoparticles clusters are no more detectable by X-ray after 2 hrs of milling, which is attributed here to their incorporation in the matrix, as opposed to dissolution in the matrix as is often suggested in the open literature. HR-TEM performed on the yttria particles before and after incorporation in the steel by ball milling shows that the yttria maintain its bcc equilibrium structure throughout the process. This could suggest that they are not dissolved during ball milling. We have clarified this point by a cross sectional examination of the EUROFER97 ball milled particles with 0.3wt% of yttria in an FIB. By X-ray EDS analysis it shows nanometric regions reach in yttrium and oxygen, which are the yttria particles embedded in the EUROFER97 particles, proving that they are not dissolved by ball milling. The mechanism at the origin of the alloy compaction during HIP was elaborated. The alloy compaction was done through HIPing at 1150°C for 2.5 hrs at a pressure of 190 MPa. This condition is shown to provide the highest bulk density, to a value close to 99.9%. In the optimization process, it is found that temperature plays a larger role than pressure for alloy compaction. The compaction appears to be mainly due to diffusion. Also, it appears that the ball milled particle size plays a major role in the alloy compaction. It is shown that smaller particles size induces better compaction, provided that the densification proceeds through diffusion flow. This further suggests to terminate the ball milling process by 20 hrs as the particles size reaches its minimum. The impact of the ball milling process on the mechanical properties was evaluated. EUROFER97 with and without ball milling, with and without oxide addition presents an average density around 99.8%. In the as received state, casted EUROFER97, atomized EUROFER97, ball milled EUROFER97 and ODS Yttria present microstructure morphology typical of tempered martensite, whereas ODS Ti presents a nano sized grain microstructure with an average grain size of 200 nm. Ball milling and HIPping does have a slight influence up to 400°C, due to the higher dislocation density induced by ball milling relative to casted EUROFER97. It becomes negligible above 400°C, since both casted and ball milled EUROFER97 present similar hardness and microstructure. In ODS Yttria, yttria presents an average particle size of 25 nm. In ODS Ti, Y-Ti-O presents an average particle size of 7 nm. In the following the identification of the hardening mechanisms in ODS F/M steels. A new model was developed to quantify the collective hardening due to different obstacles, assuming that they provide independent hardening mechanisms. The model helped identifying the dominating source of hardening in ODS F/M steels. The base of the model is the dispersion barrier strengthening model, which is given by Δσsparticle or defects = Mαμb(Nd)1/2 for the hardening due small defects or oxide particles and is given by Δσd-d = Mαμb(ρ)1/2 for the hardening due to dislocations. In this approximation, the total hardening is then given by Δσtotal = {(Δσd-d)2 + (Δσparticle or defects)2}1/2 . The model is used to explain the hardening due to oxides, dislocations and irradiation induced defects. It appears that dispersed oxides is dominating hardness up to 800°C, while following a heat treatment at 1000°C it is the dislocations that then dominate the hardness. Deformation mechanisms were investigated using differential strain rate experiments. The activation parameters, i.e. the activation volume and activation energy, were calculated. It appears that there is a decrease in the activation volume from room temperature to 400°C, which is less than 3%. At about 500°C there is a sudden drop in the activation volume that reaches a minimum at 600°C. With further increase in temperature, there is a significant raise in the activation volume. This abrupt decrease between 400°C and 600°C hints at a change in dislocation-rate controlling mechanism, marking two regimes, a low temperature one and a high temperature one. In situ TEM observations show that at low temperature the deformation is mainly due to the formation of the Orowan loops around the particle. At high temperature the activation energy is about 3.4 eV for ODS, whereas it is 1.3 eV for EUROFER97. The calculated activation energy corresponds well to the one of dislocation climb mechanism. This indicates that at high temperature the yttria particles could be overcome by climb, while at low temperature they are overcome by the Orowan mechanism. The peculiar response of the yield stress to irradiation for this class of material was rationalized using the total hardening model we developed. At a dose of 2 dpa, irradiation induced defects are higher in number density than the dispersed yttria particles, which although they are weaker than the yttria particles they can explain the significant hardening observed in the material. Conversely, the shallow impact of irradiation on ODS F/M steels at low doses, relative to base F/M steels, is explained by the fact that the amount of dispersoid is found to be higher than the density of irradiation induced defects.

EPFL2008

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

The relationship between the irradiation induced damage and the mechanical properties of single crystal Ni

Zhongwen Yao

Irradiation is known to lead to a degradation of the mechanical properties of materials. This is particularly crucial in the case of materials that will be used in the future thermonuclear fusion reactor, where extremely high irradiation doses are expected. In the quest of a better understanding between the irradiation induced defects and the mechanical properties Ni single crystal specimens have been irradiated with 590 MeV protons to doses ranging from 10-3 dpa to 0.3 dpa, at room temperature, 250°C and 350°C. The irradiation induced microstructure has been characterized by transmission electron microscopy, and the mechanical properties have been assessed by mechanical testing. Molecular dynamics (MD) simulations have been conducted in Ni and Cu in order to understand the defect formation and accumulation following a displacement cascade. The following conclusions are established. Following irradiation at room temperature, 43% to 55% of the irradiation induced defects in Ni consist in stacking fault tetrahedra (SFTs), 31%∼41% in loops and about 10 % in unidentified black dots. In the case of Ni irradiated at 250 °C, 44%∼53% of the irradiation induced defects consists in SFTs, 35%∼51% in loops, and less than 5% in black dots. In the case of Ni irradiated at 350 °C, 50% of the irradiation induced defects consists in voids. The remaining 50% include ∼ 14% SFTs and 36% loops. Moreover, it appears that these ratios are independent from the irradiation dose, contrary to what was found in the literature. These results allow clarifying a long standing issue, namely that in fact in Ni there is no transition in the ratio between SFTs and loops with increasing dose. In addition, it appears that in fact the irradiation induced defect density is similar to the one found in other irradiated fcc metals for RT irradiation. The data of the irradiation at 250°C are in agreement with previous published results on a neutron irradiation at 230 °C. It appears that the size of SFTs is independent of the irradiation dose, similar to what is found in irradiated Cu. It depends, however, on the irradiation temperature. MD simulations have been performed in order to understand the influence of the interatomic potential's parameters on the formation of defects following displacement cascades. Starting with a known defect configuration, namely the SFT, 4 different potentials are tested. Results of the formation of an SFT from a triangular platelet of vacancies show that i) the platelet of 6 vacancies did not collapse to SFT regardless of the annealing conditions, but formed a void above 800K, for all selected potentials; ii) the platelet of 15 vacancies collapsed into an SFT when simulated with Farkas-I and Farkas-II potentials, which have low stacking fault energies; iii) the platelet of 66 vacancies easily collapsed to an SFT at 500K for all applied potentials, even with a high stacking fault energy (SFE) of 300 mJ·m-2. The common neighbor analysis and Wigner-Seitz defects analysis are used for studying the resulting structures after a displacement cascade, and results show that the largest number of stacking faults is surprisingly obtained with the potential giving the highest stacking fault energy (300 mJ·m-2). An SFT-like structure appears close to the core of cascades and also an isolated interstitial loop are found with Cleri-Rosato potential. It appears that the defects size, configuration and density relates more to the displacement threshold energy than to the stacking fault energy. It appears that a significant irradiation hardening occurs starting at the lowest dose of 10-3 dpa, at room temperature and 250°C. With increasing dose, the yield shear stress increases. The dose dependence of this increase varies from an irradiation temperature of room temperature to 250°C. In fact the irradiation hardening shows a significant temperature dependence. The thermal activation energy was calculated by measuring the activation volume, which is obtained from relaxation tests. In unirradiated Ni, an activation energy of about 0.3 eV can be extrapolated at RT. In case of irradiated Ni, the unsuccessful fitting with a linear relationship between the thermal energy and temperature suggests that multiple deformation mechanisms are operating simultaneously or in turn over the considered temperature range of –196 °C to 423 °C. The dispersion barrier hardening model is used to interpret the irradiation induced hardening in irradiated Ni. According to the equation, Δτ = αμb(Nd)1/2, with α, the obstacle strength. It appears that α is 0.12 for the room temperature irradiation induced hardening, similar to what was previously found for irradiated Cu, and is 0.22 for the hardening following irradiation at 250°C, which is much larger than previous values. This difference in α is suggested to be related to the presence of voids at high temperatures. It is suggested that voids are stronger obstacles than SFTs or dislocation loops. In situ TEM observations were made in irradiated Cu, in order to assess the strength of irradiation induced defects as obstacles to gliding dislocations. It appeared that the strength of obstacles is approximately 100 MPa. Considering that 90% of the damage consists in SFTs, it is suggested that the obstacles presenting a resistance of 100 MPa are SFTs. In particular, the pulling out of a string from the interaction between what was assumed to be an SFT and a moving dislocation could be recorded for the first time. This reaction implies that the gliding dislocation is strongly pinned by SFTs. TEM observation of the microstructure of the irradiated Ni following plastic deformation shows that defect-free channel formation is the deformation mechanism at the beginning of yielding. The channeling explains the observed softening. As deformation proceeds, a dislocation cell structure is initiated. It takes different forms depending on whether the sample is irradiated or not. In the irradiated case, cell formation initiates from mobile dislocation in the defect free channels. This implies that the eventual cell configuration will retain part of the dimensionality of the channels, for the least their width of 100 nm for the cell size, because of the geometrical constraint. In the unirradiated case the dislocation cell size is larger. The higher flow stress observed in the irradiated case is thus related to the dislocation cell size.

EPFL2005