EBSD coupled to SEM in situ annealing for assessing recrystallization and grain growth mechanisms in pure tantalum

An in situ annealing stage has been developed in-house and integrated in the chamber of a Scanning Electron Microscope equipped with an Electron BackScattered Diffraction system. Based on the Joule effect, this device can reach the temperature of 1200°C at heating rates up to 100°C/s, avoiding microstructural evolutions during heating. A high-purity tantalum deformed sample has been annealed at variable temperature in the range 750°C-1030°C, and classical mechanisms of microstructural evolutions such as recrystallization and grain coarsening phenomena have been observed. Quantitative measurements of grain growth rates provide an estimate of the mean grain boundary mobility, which is consistent with the value estimated from physical parameters reported for that material. In situ annealing therefore appears to be suited for complementing bulk measurements at relatively high temperatures, in the context of recrystallization and grain growth in such a single-phase material. © 2013 The Authors Journal of Microscopy © 2013 Royal Microscopical Society.

Finite element modeling and experimental study of brittle fracture in tempered martensitic steels for thermonuclear fusion applications

Pablo Federico Mueller

In this work we have studied brittle fracture in high-chromium reduced activation tempered martensitic steels foreseen as structural materials for thermonuclear fusion reactors. Developing the adequate materials that can withstand the severe irradiation conditions of the burning plasma in a future fusion reactor is one of the major challenges to be solved in order to make profit from the great advantages of thermonuclear fusion as an energy source. High-chromium tempered martensitic steels such as F82H and the most advanced version Eurofer97 are among the main candidate materials for structural applications in future fusion power plants due to low irradiation-induced swelling, good mechanical and thermal properties, and reasonably fast radioactive decay. The most concerning drawback of these kind of steels is irradiation embrittlement, which is manifested by a ductile-to-brittle transition temperature shift to higher temperatures after irradiation whose amplitude depends on the irradiation conditions (temperature, neutron flux, neutron fluence, etc). The aim of this work was to study and model brittle fracture in the ductile-brittle transition region of this kind of steels in the as-received unirradiated conditions. It is necessary to be able to transfer laboratory specimen fracture data to real components and structures in order to assess the performance of these steels in the different operating and transient conditions they could find during the operation life of a fusion reactor. In order to do so, the specimen geometry effects and specimen size effects on measured fracture toughness need to be properly understood, taken into account and predicted with an appropriate model. In particular, specimen size effect on measured toughness is a major concern for the nuclear materials research community owing to the limited irradiation volume in current and planed materials irradiation facilities. The main results of this PhD work are summarized below. The microstructure of Eurofer97 and F82H has been characterized and compared by means of optical microscopy, scanning electron microscopy, transmission electron microscopy and energy-dispersive X-ray spectroscopy in order to identify microstructural features that could play a role in the measured fracture toughness. Both steels have similar but slightly different chemical composition and final heat-treatments but the prior austenitic grain size measured in F82H is approximately 8 times larger than in Eurofer97. It was shown that the alloying element Tantalum, added to stabilize the austenite grain size, played a different role in both steels. After a careful analysis of the particles present in both steels, it was found that Tantalum in Eurofer97 formed carbides of an average size around 100 nanometers. In contrast in F82H it did not form small carbides but formed big oxide inclusions with a size up to 30 µm. These large particles do not effectively pin the grain boundaries. The different behavior of Tantalum in these steels is believed to be mainly a consequence of the larger content of Oxygen present and the smaller amount of Aluminum in F82H compared to Eurofer97. The Master-Curve ASTM-E1921 standard is a method initially developed to determine the ductile-to-brittle transition reference temperature, T0, in fission reactor ferritic steels from a small number of experiments. In this work the applicability of the Master-Curve method to reduced activation tempered martensitic steels such as Eurofer97 and F82H was studied in detail. Fracture tests with pre-cracked sub-sized compact tension specimens (three different sizes, 0.18T, 0.35T and 0.87T of Eurofer97) were carried out in the temperature range [-196 °C, -40 °C]. The toughness-temperature behavior and scatter were shown to deviate from the ASTM-E1921 standard predictions near the lower shelf. Using the method of maximum likelihood, the athermal component of the Master-Curve was calculated to better fit our fracture toughness data from the lower to the middle transition region. We showed that these Master-Curve adjustments are necessary to make the T0 values obtained near the lower shelf with 0.35T size C(T) specimens consistent with those obtained in the middle transition region with 0.87T C(T) specimens. The ASTM-E1921 specimen size limitations, setting the maximum toughness measurable with a given specimen size, were found to be too lenient for this kind of steels. This problem was especially evident in fracture toughness data of Eurofer97 and F82H obtained in the upper transition temperature range with two different specimen sizes. Thus a more stringent specimen size requirement was proposed to avoid inconsistent transition temperature determinations. A promising local fracture model with the potential of predicting cleavage fracture toughness was studied. Finite element simulations were undertaken for compact specimens, notched specimens and tensile specimens of Eurofer97 steel tested from 20 °C down to -197 °C. Three and two dimensional as well as axisymmetric simulations were run in order to calculate the stress and strain fields at the onset of brittle fracture. For each tested temperature, the calculated load-displacement curves were found to reproduce very well those of the experiments. A local approach fracture criterion was studied. This criterion states that when the maximum principal stress is larger than a critical stress within a critical volume ahead of the crack tip, notch or neck, cleavage is triggered leading to macroscopic fracture of the specimen. It was shown that this model is able to predict the minimum fracture load of the notched specimens with the same values of critical stress and critical volume that were calibrated to predict the lower bound of fracture in compact specimens. This local approach model was also successfully used to predict the strong size effect observed experimentally in pre-cracked compact tension specimens in the upper transition region. The critical fracture stress determined in the standard tensile specimens was found higher than that of the fracture specimens. It was suggested that this difference stems from a significant difference in the stress state between the different specimens (triaxiality). Finally, the comparison between the overall fracture behavior in the transition between F82H and Eurofer97 steels indicated that these two materials are quite similar. A difference in the reference temperature T0 of about 20 °C was found, with a nominal value of -100 °C and -80 °C for F82H and Eurofer97 respectively. However, some excessive scatter in the toughness data was found in F82H with more data points than expected lying below the 1% tolerance bound; this was not observed in Eurofer97. From an engineering point of view, the entire fracture databases of these two steels are encompassed practically by the same lower tolerance bound.

EPFL2009

Microstructure and embrittlement of leaded copper alloys

Laurent Felberbaum

Mechanisms responsible for the embrittlement of leaded "free-machining" copper alloys at intermediate temperatures (i.e. around 400 °C) are examined using a combination of microstructural and mechanical characterization. Tensile tests are conducted on pure copper and on industrial high-strength Cu-Ni-Sn alloys, comparing leaded and unleaded versions, at strain rates ranging from 10-4 to 10 s-1, and various temperatures between 25 °C and 800 °C. It is found that, in the leaded samples, both liquid metal embrittlement and grain boundary embrittlement operate at intermediate temperature,i.e., near the melting point of lead. Their respective importance depends mainly on the strain rate: fracture at high strain rate (10 s-1) displays characteristic signatures of liquid metal embrittlement, while at low strain rate (10-4 s-1) the ductility trough behaviour can be attributed to grain boundary embrittlement. In leaded copper and copper alloys, lead is mostly present as discrete inclusions, frequently located along grain boundaries. The dihedral angle of intergranular lead inclusions is then the main microstructural parameter besides their size. A new technique is proposed and demonstrated for the measurement of inclusion dihedral angle in a high-purity Cu-1 wt.% Pb alloy, based on room temperature characterization of samples that were held at elevated temperature at times sufficiently long for shape equilibration of the inclusions and then rapidly quenched. Quantitative scanning electron microscopic analysis of metallographic surfaces along which individual inclusions have been dissolved is used to deduce the dihedral angle using a mathematical fit of their solid/liquid interface. For a specific temperature, we show that the dihedral angle is not unique; this reflects the fact that high-angle grain boundary energies are not constant in copper. Existing micromechanical analyses of shape equilibrium for intergranular inclusions in the presence of external stress are adapted to produce an improved description at low stress and to take into account the inclusion bulk compressibility. The derivation is based on minimization of the global capillary and elastic strain energy associated with such inclusions. An adimensional parameter that depends on the applied stress, the inclusion volume, the elastic properties of the solid, and the relevant interfacial energies emerges from the analysis as the sole factor governing the shape of the inclusion. If void nucleation occurs easily within the inclusion, such that its apparent bulk modulus is nil, the inclusion become unstable, degenerating to a crack when this adimensional parameter exceeds a critical value. Measurements of the dihedral angle on samples previously subjected to a remote stress at 400 °C show that applied stress causes a reduction in the apparent dihedral angle of liquid lead inclusions. Calculated critical stresses for both leaded pure copper and leaded copper-nickel-tin alloys that fail at high strain rate by liquid metal embrittlement are calculated for individual inclusions in the metal. These calculated critical stresses correlate relatively well with the measured fracture stress, suggesting that shape instability of the inclusions cause fracture of the alloys. This result is applied towards the suggestion of general strategies for reduction of the susceptibility of leaded copper alloys to liquid metal embrittlement. A case study addressing the problem of quench cracking of leaded Cu-Ni-Sn alloys on the basis of experimentation and finite element thermomechanical simulation, showing how the problem can be alleviated in production, concludes the thesis.

EPFL2005

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