Soudage laser or-acier | EPFL Graph Search

In jewelry manufacturing, joining dissimilar materials is usually achieved by brazing or soldering, in which the materials comprising the joint are heated to a suitable temperature in the presence of a filler metal having a liquidus below the solidus of the base metals. This simple and cost-effective method has however several disadvantages, including undesired phases transformations in the bulk base metals during heat treatment. An alternative joining methodology is laser beam welding. Due to its high energy density, a laser beam is able to melt locally the interface, without affecting the bulk microstructure of the parts to be joined. With this method, however, it is necessary to have a thorough understanding of microstructure formation during solidification. In this thesis, the process of laser welding an 18 carat yellow gold to a superaustenitic stainless steel has been studied and developed. In order to understand the microstructure formation during this welding process, a fundamental metallurgical analysis of the gold-steel system was first carried out, partially based on the binary Au-Fe elemental system. Following this, laser welds have been characterized and a concept for gold-steel welding proposed and successfully tested. First, Differential Thermal Analysis (DTA) and directional solidification in a Bridgman-type furnace followed by rapid quenching were used to acquire an understanding of the microstructure evolution upon cooling in Au-Fe alloys. During the (γ-Fe) primary phase solidification, micro- and macrosegregation phenomena influence the final composition profile. The peritectic phase forms rapidly and instabilities develop during solid-phase peritectic transformation. Below the peritectic temperature, solid phases decompose in a miscibility gap, leading to continuous or discontinuous precipitation, depending on cooling conditions. Finally, (α-Fe) ferrite precipitation gives insights on orientation relationships between primary and peritectic phases. Second, yellow gold-stainless steel alloys have been studied at low cooling rates by DTA. By combination of these observations with the Au-Fe metallurgical analysis, the microstructure formation in gold-steel laser welding could be explained. It appears that the microstructure formed during laser welding is strongly related to the local melt pool composition. Three different cases were identified: (1) Due to its small solidification interval and little segregation, goldrich liquid solidifies quickly as columnar grains, avoiding formation of cracks or porosities. (2) Due to the high cooling rate, liquid of intermediate composition is brought rapidly to low temperature, reaching the metastable liquid miscibility gap, resulting in demixing into two liquids. (3) A liquid containing mainly iron will solidify with a dendritic microstructure rejecting large amount of gold in the interdendritic liquid. In this last case, this large solidification interval combined with the considerable stresses around the weld pool leads certainly to cracking. Based on the above analysis, a concept of gold-steel laser welding can be proposed. By focusing the laser beam on the gold side of the interface, a growing liquid pool can be formed in this metal. When the melt pool reaches the steel surface, a layer of peritectic phase will form, without melting much steel. Since the liquid pool contains mainly gold, hot cracking is avoided. After solidification, the peritectic layer ensures a smooth transition between the mechanical properties of gold and steel. This concept has been successfully applied on a number of real pieces, forming sound welds with suitable properties. This thesis has also shown that a fundamental metallurgical approach is a powerful tool for complex welding development. It allows not only to develop a viable industrial solution, but also to acquire a knowledge that can be used to expand the process to other materials or different geometries.

Peritectic solidification of Cu-Sn alloys

Frédéric Kohler

Many commercially important alloys such as Fe-C, Fe-Ni, Cu-Zn, Cu-Sn and Ti-Al exhibit peritectic transitions on solidification. In the two-solid phase region of peritectic metallic systems, a wide spectrum of microstructures has been revealed during directional solidification experiments at low growth rates, where both α- and β-phases would grow independently as planar fronts [1, 2, 12–20]: discrete bands, islands, simultaneous growth of the α- and β-phases in the form of oriented lamellae or fibers and continuous oscillatory tree-like microstructures in highly convective peritectic systems [3, 21, 22]. Recently, Vandyoussefi et al. [27] and Dobler et al. [16] have done extensive solidification experiments on Fe-Ni alloys in the corresponding two-solid phase region. Depending on the growth conditions and local composition, a lamellar eutectic-like structure was observed. Despite recent works, the peritectic solidification at low growth rate is not fully understood. Additionally, coupled growth was observed only for peritectic systems for which the solidification interval of the primary phase ΔΤ 0α is fairly small and with negligible convection; but is it also possible for peritectic alloys with fairly large ΔΤ 0α such as Cu-Sn? If so, what would be the effect of convection? And finally, is the coupled growth front really isothermal? In the present contribution, the peritectic Cu-Sn system has been chosen because of its remarkable properties and technological importance. Indeed, two essential features distinguish Cu-Sn alloys from other peritectic systems: The equilibrium solidification interval of the a-phase decreases with the tin concentration in the hypoperitectic composition range. Additionally, the equilibrium solidification interval of the primary phase is about 25 times larger in the Cu-Sn system than in the Fe-Ni peritectic system. Since only few studies have been made on the detailed solidification of tin bronze alloys, the first goal of this study was to carry out thermal analyses in order to investigate precisely the solidification of Cu-Sn alloys and to measure temperatures of solid-state phase transformations. Therefore, an SPTA assembly has been built and successfully calibrated. Additionally, a heat flow model was developed in this work and coupled with a Cu-Sn thermodynamic database to treat solidification-dependent latent heat. Single pan thermal analyses on three Cu-Sn alloy compositions revealed that the corresponding phase diagram recently re-investigated in Liu et al. review [28] is the most reliable. Directional solidification experiments on Cu-Sn alloys of various compositions have been conducted at different velocities in a high gradient Bridgman furnace. In this context, the set-up already used by Dobler [14] has been modified in order to reduce the size of the samples and, accordingly, natural convection and the associated macrosegregation. During each solidification run, two different geometries were thus tested (3 mm diameter cylinder and 4/6 mm tube), allowing the observation of convection effects on solidifying microstructures. An essential outcome of this project is that α+β lamellar structures have been observed in both sample geometries, even when Vp is above the critical velocity for the α planar front growth, i.e., the α-phase grew with a cellular morphology. With the Cu-Sn phase diagram used in the present work and for positions close to the quenched interface, concentration profiles measured across the lamellar structures revealed that the corresponding operating temperature Τ*(x) seems to be non-isothermal, in contradiction with previous experimental observations in Fe-Ni peritectic alloys [14–16]. But, since the Cu-Sn phase diagram is not assessed accurately, Τ*(x) is still subject to large variations (±10°C). On the other hand, the peritectic transformation after solidification has been evidenced and modifies the compositions and fractions of phases. Additionally, local concentration showed that a partial banding mechanism and an unstable lateral propagation of both solid phases were responsible for the formation of lamellar structures. Various geometrical arrangements were revealed on transverse cuts of lamellar structures: straight α- and β-lamellae, labyrinth-like microstructures, disordered arrays of rods of one phase in the matrix of the other phase and uniform regions of β-phase. The dynamic evolution of lamellar structures was attributed to growth competition mechanisms and instabilities between phases originated from the same nuclei. Finally, it should be noted that the Kurdjumov-Sachs relationship was observed between the primary and the peritectic phases on both the straight lamellae and labyrinth-like microstructure. Finally, all lamellar structures ended by the termination of α-lamellae and the subsequent growth of a β planar front. In this context, the α-lamellae of lamellar structures observed in the 4/6 mm tube exhibited in all cases a "hammer-like" morphology before being overgrown by the peritectic phase. Careful investigations of this morphology revealed that it resulted effectively from growth competition mechanisms, which remain to be fully understood. In a simulated geometry corresponding to the 4/6 mm tube of Cu-Sn samples, 3D macrosegregation modeling revealed that a nearly helical convection pattern existed in the liquid phase ahead of a moving planar α-liquid interface. And between each plume, the fluid flow was mostly oriented parallel to the solid-liquid interface. As a consequence, the lateral propagation of the α- and β-phases during the 3D growth competition was surely affected by these convective motions. Finally, natural convection and the associated solute enrichment were assumed to be responsible for the destabilization of lamellar structures and the subsequent end of α-lamellae. The multi-phase field computations emphasized that the stability of the triple junction between α-, β- and liquid phases is function of the initial alloy concentration C0, of the lamellar spacing λ and of the thermal gradient Gl. In a near future, directional solidification experiments are planned to be conducted onboard of the International Space Station. Then, the obtained peritectic microstructures will have to be compared with the microstructures observed in the present work in order to assess more precisely the effect of natural convection on the formation of lamellar structures.

EPFL2008

Analysis of nano-sized irradiation-induced defects in Fe-base materials by means of small angle neutron scattering and molecular dynamics simulations

Gang Yu

Thermonuclear fusion of light atoms is considered since decades as an unlimited, safe and reliable source of energy that could eventually replace classical sources based on fossile fuel or nuclear fuel. Fusion reactor technology and materials studies are important parts of the fusion energy development program. For the time being, the most promising materials for structural applications in the future fusion power reactors are the Reduced Activation Ferritic/Martensitic (RAFM) steels for which the greatest technology maturity has been achieved, i.e., qualified fabrication routes, welding technology and a general industrial experience are almost available. The most important issues concerning the future use of RAFM steels in fusion power reactors are derived from their irradiation by 14 MeV neutrons that are the product, together with 3.5 MeV helium ions, of the envisaged fusion reactions between deuterium and tritium nuclei. Indeed, exposure of metallic materials to intense fluxes of 14 MeV neutrons will result in the formation of severe displacement damage (about 20-30 dpa per year) and high amounts of helium, which are at the origin of significant changes in the physical and mechanical of materials, such as hardening and embrittlement effects, for instance. This PhD Thesis work was aimed at investigating how far the Small Angle Neutron Scattering (SANS) technique could be used for detecting and characterizing nano-sized irradiation-induced defects in RAFM steels. Indeed, the resolution limit of Transmission Electron Microscopy (TEM) is about 1 nm in weak beam TEM imaging, and it is usually thought that a large number of irradiation-induced effects have a size below 1 nm in RAFM steels and that these very small defects actually contribute to the irradiation-induced hardening and embrittlement of RAFM steels occurring at irradiation temperatures below about 400°C. The aim of this work was achieved by combing SANS experiments on unirradiated and irradiated specimens of RAFM steels with Molecular Dynamics (MD) simulations of main expected nano-sized defects in irradiated pure Fe and Fe-He alloys, as model materials for RAFM steels, and simulations of their corresponding TEM images and SANS signals. In particular, the SANS signal of various types of defects was simulated for the first time. The methodology used in this work was the following: SANS experiments were performed by applying a strong saturating magnetic field to unirradiated and irradiated specimens of three types of RAFM steels, namely the European EUROFER 97, the Japanese F82H and the Swiss OPTIMAX A steels. The available irradiated specimens included specimens which had been irradiated with 590 MeV protons in the Proton IRradiation EXperiment (PIREX) facility at the Paul Scherrer Institute (PSI) at temperatures in the range of 50-350°C to doses in the range of 0.3-2.0 dpa. SANS spectra as well as values of the so-called A ratio, which represents the ratio of the total scattered intensity to the nuclear scattered intensity, were obtained for the various irradiation doses and temperatures investigated. MD simulations of atomic displacement cascades in pure Fe and in Fe-He alloys were performed using Embedded Atom Method (EAM) many-body interatomic potentials. The main nano-sized defects that should be produced in RAFM steels under irradiation were created by means of MD in pure Fe. These included dislocation loops of various types, voids, helium bubbles with various He concentration and Cr precipitates. TEM images of cascade damage and all the defects created by MD were simulated in the dark field/weak beam imaging modes by using the Electron Microscopy Software (EMS) developed by P.A. Stadelmann (EPFL) and analyzed in terms of variations of contrast intensities versus depth inside the specimen. The SANS signal provided by cascade damage and all the defects created by MD was simulated by using a slightly modified version of EMS, accounting for neutrons instead of electrons. The SANS technique has been proven in this work to be a very powerful tool for detecting nano-sized irradiation-induced defects and a tool well complementary to TEM for characterizing such very small irradiation-induced defects. Indeed, TEM appears most adapted to investigate structural defects, such as dislocation loops and helium bubbles with high helium concentration, which yield significant lattice deformation of the surrounding matrix, while SANS is most adapted to investigate phase defects, such as voids, helium bubbles with low helium concentration and Cr precipitates. By combining the results of SANS experiments with those of MD simulations, TEM image simulations and SANS signal simulations, the nano-sized irradiation-induced defects were tentatively identified as small helium bubbles. While the radiation hardening measured for RAFM steels cannot be explained by accounting only for the defects observed in TEM, it could be successfully modeled by accounting also for a reasonable number density of the nano-sized defects evidenced using the SANS technique.

EPFL2008

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

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

Amuthan Ramar

EPFL2008

Analysis of nano-sized irradiation-induced defects in Fe-base materials by means of small angle neutron scattering and molecular dynamics simulations

Gang Yu

EPFL2008

Peritectic solidification of Cu-Sn alloys

Frédéric Kohler

EPFL2008