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Slagsegheten hos stål är komplext beroende av ett antal olika parametrar, som i olika grad påverkar varandra. Många betydande och viktiga ledtrådar har hittats som förhoppningsvis kan ge en bredare förståelse varför slagsegheten varierar och varför den inte alltid uppnås i nuvarande process. Studien har utförts på kalldragna sömlösa rör, med stålsort EN 10305-1. Projektet innefattar ett antal olika produkt specifikationer med olika dimensioner, med lite olika kemiska sammansättning. På grund av slagseghetskraven på industriella produkter var det viktigt att utvärdera vilka parametrar som påverkar. En utvärdering gjordes dels av resultat från slagseghetsprovning efter påverkande processer, och dels från utvärdering av stålets mikrostruktur. Förutom dimension och kemisk sammansättning, varierar processlinjens faktorer. Den första processen som påverkar är kalldragningen som i denna studie har haft reduktioner mellan 15% och 31%. Avspänningsglödgningens temperatur påverkar även, där temperaturerna 525 oC, 575 oC och 600 oC är de som använts på produkterna i projektet. Vissa av projektets produkter har även normaliserats, och resultaten tyder på att detta är en mycket viktig process om man ser till slagseghet. Målsättningen var att får fram omslagskurvor vid olika tillstånd av stålet. Utifrån dessa tillsammans med utvärdering av stålets mikrostruktur och kemisk sammansättning kunde flera diskussioner föras. Den viktiga slutledningen som kunde dras var att inga faktorer är ensam avgörande för slagsegheten i den här studien. Den kemiska sammansättningen påverkar ihop med processlinjens olika faktorer. Man bör ligga på en temperatur för avspänningsglödgning över 550 oC. Projektet visar även på att större dimensioner påverkar slagsegheten negativt. Två viktiga felmarginaler kan ha påverkat resultaten i projektet, dels positionen för vart anvisningen görs på provstaven i förhållande till rören och dels användandet av en praxis som beräknar resultat från provning med reducerade provstavar. En undersökning av dessa felmarginaler gjordes, och resultaten var tydliga. En mer noggrann provuttagning bör göras, och man bör vara mer försiktig vid användandet av praxisen.

<p>Rapid dissolution of alumina upon addition to the cryolitic melt is crucial for the modern Hall-Heroult process for aluminium production. The formation of slow - dissolving alumina agglomerates may be detrimental, and irregular dissolution kinetics may cause the loss of process control. So-called anode effects may subsequently ignite, which are a major source of green-house gases from the primary aluminium industry.</p><p>A literature review and the study of the theory of sintering provides the background for discussing the present work. The most probable mass transport mechanism in the transition alumina-fluoride-moisture system studied here is surface diffusion. Surface diffusion is a non-densifying mass transport mechanism that will result in coarsening (alumina grain growth) but only weak interparticle bonding since no macroscopic shrinkage is involved. Rapid mass transport is known to result when there is a simultaneous phase transformation, and this is the case when transition alumina transforms to α-alumina, catalyzed by the presence of fluorides.</p><p>The main experimental techniques used in the present work were powder X-ray diffraction (XRD) and Scanning Electron Microscopy (SEM). Supporting techniques used have been speciffc area determination by the BET method and simple thermo-gravimetric techniques. An optical furnace was designed and built in order to study the dissolution of tablet alumina agglomerates.</p><p>A preliminary agglomeration study of preformed cylindrical alumina samples served to map some of the most important mechanisms involved when alumina powder interacts with alumina-saturated cryolitic melt. The conditions at the alumina-melt interface were studied, but it is concluded that the experimental method could not provide the necessary parameter control in order to study the agglomeration mechanism in further detail.</p><p>The tablet agglomerate study is the major experimental contribution of the present work. The experimental method provided good control of the sample chemistry and well defined temperature and time variables. It is concluded that liquid cryolitic melt (NaAlF₄) provides an effective mass transport route for the transformation assisted growth of α-alumina platelets. The platelets that initially form will provide the limited mechanical strength necessary for agglomerate formation and their persistence in a cryolitic melt. Alumina agglomeration may therefore take place with only partial, initial phase transformation. It is concluded that differences in the agglomeration behavior of various qualities of alumina may be the rate determining property for alumina dissolution kinetics in cryolitic melts. Differences in the agglomeration behavior may be due to a number of physical properties of alumina. It is argued here that the fundamental, but difficult to measure, alumina nano-structure may be most important.</p><p>The alumina nano-structure is correlated to secondary alumina properties such as the α-alumina content, specific surface area (BET) and moisture content (MOI, LOI). In this study an X-ray diffraction line profile analysis using the Warren-Averbach method shows that there is a significant difference in the nano-structure of the two smelter grade alumina qualities under study. This may explain the different agglomeration behavior that is observed.</p><p>An optical study of tablet agglomerate dissolution in cryolitic melt proved to be largely unsuccessful due to severe corrosion of the quartz crucibles used. However, a proposed mechanism for the tendency for disintegration of alumina agglomerates, thus dissolving as \snow-flakes" is supported. </p><p>The temperature response time in the tablet alumina samples was studied in order to determine the experimental limit of the shortest time period possible in the experiments. The exothermal γ -> α transformation is observed for secondary alumina samples containing adsorbed fluorides. An interesting effect of the carbon content in secondary alumina is also shown.</p><p>The moisture content of smelter grade alumina is a function of the alumina quality, in particular the technology used for the calcination of the aluminium trihydrate precursor. In the current study the moisture content is shown to be a dynamic function of the ambient temperature and relative humidity. The moisture content is an important variable for the study of alumina agglomeration, and for the fluoride emission from the Hall-Heroult process. The kinetics of moisture desorption and absorption for various alumina qualities is studied. The desorption kinetics is concluded to be signifcantly different, while it is also shown that practical absorption kinetics is a function of the sample size and available surface area. </p>

<p>A comprehensive study of the softening behaviour of two different non-heat treatable AlMn-alloys has been carried out. These alloys were a laboratory processed and an industrially processed AA3103-alloy. The primary objective of the laboratory processed alloy has been directed towards the relationship between the amount of manganese in supersaturated solid solution and the material behaviour during deformation and annealing. The focus for the industrially processed alloy was a detailed material characterisation during hot and cold rolling as well as to follow the softening behaviour of the alloy.</p><p>The cold rolled microstructures were characterised with respect to subgrain size, misorientation across the subgrain boundaries, particle break-up and global texture. As far as it concerns the cold deformed microstructure it was found that the different homogenisation treatments, resulting in different amount of Mn in supersaturation, had little effect on the deformed microstructure.</p><p>Detailed experimental work on the softening behaviour after cold deformation and the interaction between recrystallisation and precipitation, i.e. concurrent precipitation, has been carried out. Hardness and yield stress measurements, which defines the strength of the material, revealed that the softening behaviour was significantly slowed down in case of concurrent precipitation. It was further found that the precipitation reaction in this case occurred on the subgrain boundaries. Thus the precipitates considerably retarded the recrystallisation reaction as compared to the case when no precipitation occurred.</p><p>TTT-diagrams have been constructed on the basis of hardness and conductivity measurements. From these diagrams a characteristic temperature, T_c , for the different material conditions are identified. It was found that as the annealing was carried above this temperature the microstructure consisted of a fine grained equiaxed microstructure. Below T_c the grains become elongated in the rolling direction and the average grain size became much coarser.</p><p>With respect to recrystallisation texture, a very strong P-texture and in addition some ND-rotated cube texture was found in case of concurrent precipitation. This viistrong P-texture was investigated in detail, and it can be concluded that particle stimulated nucleation of recrystallisation (PSN) plays a significant role in the nucleation of these components. It was found that these texture components had a growth advantage in the early stage of annealing and that they are most probably a result of microgrowth selection, which often are related to a nucleation effect. When precipitation did not interact with recovery and recrystallisation the recrystallisation texture was either random or consisted of a weak cube texture.</p><p>The softening behaviour of the different materials has been modelled by a physically based softening model, which involves both the recovery and recrystallisation reactions. Both for the AlMn-alloys and for an additional commercially pure Al-alloy, relatively good model predictions were obtained for the softening behaviour when recrystallisation occurred prior to precipitation. However, when the softening reactions were retarded by heavy precipitation of dispersoids the model, in its original form, was not capable to predict the softening behaviour. In this case an additional retarding drag, which reduces the total number of viable recrystallisation nuclei, was added. With the addition of this drag relatively good model predictions were also obtained in case of concurrent precipitation. However, the model was not capable to predict the softening behaviour of the industrially processed AlMn-alloy particularly well.</p>

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<p>The overall objective of this work has been to develop aluminium alloys, which after hot and cold deformation are able to withstand high temperatures without recrystallising. This has been done by investigating aluminium alloys with various additions of hafnium, scandium and zirconium, with a main focus on Hf and to which extent it may partly substitute or replace Zr and/or Sc as a dispersoid forming elements in these alloys.</p><p><i>What is the effect of hafnium, alone and in combination with Zr and/or Sc and how do hafnium containing alloys perform?</i></p><p>It is shown that hafnium may alter or modify the casting structure, though, not to the better as it can form TCGs in combinations with Zr and/or Sc. This is not advantageous neither as far as it concerns grain refining nor precipitation of dispersoids.</p><p>When precipitation of binary Al-Hf is compared to Al-(Hf)-(Zr) alloys, hafnium shows even slower precipitation than in Al-(Hf)-(Zr) alloys and also much slower and with a poorer spatial distribution of dispersoids than in Al-Sc or Al-Hf-Sc alloys. As a consequence, it may be concluded that binary aluminium-hafnium alloys are of limited interest as they display a poor recrystallisation resistance when no other alloying elements are added.</p><p>However, when hafnium is added together with scandium and/or zirconium, precipitation may actually improve both in Al-Hf-Zr alloys and in scandium containing alloys like Al-Hf-Sc and Al-Hf-Sc-Zr. Hafnium can still not completely replace neither Zr nor Sc due to its poor precipitation properties. However, Al-Hf-Sc-Zr alloys show a better dispersoid distribution (number density and volume fraction) than what was observed in Al-Sc-Zr alloys after extrusion. The Al-Hf-Sc-Zr alloys also show extreme high temperature properties, by withstanding recrystallisation at high temperatures and long annealing times. After severe cold deformation, the extruded profiles of both the Al-Sc-Zr alloy and the Al-Hf-Sc-Zr alloy displayed a remarkable recrystallisation resistance.</p><p>The reason why the Al-Zr-Sc- and the Al-Hf-Zr-Sc alloys behave so well has been investigated by detailed 3D Atom Probe investigations. Due to the homogeneous precipitation of Al3Sc dispersoids and the retarding effect from Hf and/or Zr containing shells, thus limiting the coarsening of these dispersoids, these combinations have been shown successful.</p><p>As we have seen when Hf and Zr are added in combination equal or improved recrystallisation properties can be obtained. This is specially the case when both these elements are added together with scandium. Since Hf and Zr are extremely difficult (and thus costly) to separate, Al-Zr master alloys used in industry today can in the future probably contain more Hf, lowering the cost of master alloys.</p><p>The work presented in this thesis have hopefully added some new insight and a better understanding of the effects of adding various dispersoid forming elements to aluminium, alone and in various combinations, which may be useful for industry today and a basis for further alloy development.</p>

<p>Cold chamber high pressure die casting, (HPDC), is an important commercial process for the production of complex near net shape aluminium and magnesium alloy castings. The work presented in the thesis was aimed at investigating the microstructure formation in this type of casting. The solidification characteristics related to the process and the alloys control the formation of grains and defects. This again has a significant impact on the mechanical properties of the castings.</p><p>The investigations were carried out mainly using the AM60 magnesium alloy and the A356 aluminium alloy. Two different casting arrangements were used: the cold chamber HPDC and the gravity die casting methods, which allowed for different flow and solidification conditions. The microstructures in the castings were investigated using optical microscopy, image analysis, scanning electron microscopy, electron back scatter diffraction measurements and electron probe microanalysis.</p><p>In the HPDC experiments, the shot sleeve solidification conditions were investigated primarily by changing the melt superheat on pouring. This significantly affected the microstructures in the castings. The fraction of externally solidified crystals (ESCs) was consistently found to be largest near the gate in both the AM60 and the A356 die castings. This was attributed to the inherent shot sleeve solidification conditions and the flow set up by the plunger movement. When the superheat was increased, a lower fraction of ESCs was found in the castings. Furthermore, a high superheat gave ESCs with branched dendritic/elongated trunk morphology whilst a low superheat generated coarser and more globular ESCs, both in the AM60 and the A356 castings. The ESCs typically segregated towards the central region of the cross sections at further distances from the gate in the die castings.</p><p>When a thin layer of thermal insulating coating was applied on the shot sleeve wall in the production of AM60 die castings, it nearly removed all ESCs in the castings. Using an A356 alloy, (and no shot sleeve coating), with no Ti in solution gave a significantly lower fraction of ESCs, whereas AlTi5B1 grain refiner additions induced an increase in the fraction of ESCs and a significantly finer grain size in the castings. The formation of globular ESCs was enhanced when AlTi5B1 grain refiner was added to the A356 alloy.</p><p>In controlled laboratory gravity die casting experiments, typical HPDC microstructures were created by pouring semi-solid metal into a steel die: The ESCs were found to segregate/migrate to the central region during flow, until a maximum packing, (fraction of ESCs of ~35-40%), was reached. The extent of segregation is determined by the fraction of ESCs, and the die temperature affects the position of the ESCs. The segregation of ESCs was explained to occur during flow as a result of lift forces.</p><p>The formation of banded defects has also been studied: the position of the bands was affected by the die temperature and the fraction of ESCs. Based on the nature of the bands and their occurrence, a new theory on the formation of defect bands was proposed: During flow the solid distribution from the die wall consists of three regions: 1) a solid fraction gradient at the wall; 2) a low solid fraction region which carries (3) a network of ESCs. A critical fraction solid exists where the deformation rate exceeds the interdendritic flow rate. When the induced stress exceeds the network strength, deformation can occur by slip, followed by liquid flow. The liquid flow is caused by solidification shrinkage, hydrostatic pressure on the interior ESC network, and gaps forming which draw in liquid.</p>

<p>This doctoral thesis is concerned with rapid solidification of AB₅ materials suitable for electrochemical hydrogen storage. The primary objective of the work has been to characterise the microstructure and crystal structure of the produced AB₅ materials as a function of the process parameters, e.g. the cooling rate during rapid solidification, the determination of which has been paid special attention to.</p><p>The thesis is divided in to 6 parts, of which Part I is a literature review, starting with a short presentation of energy storage alternatives. Then a general review of metal hydrides and their utilisation as energy carriers is presented. This part also includes more detailed descriptions of the crystal structure, the chemical composition and the hydrogen storage properties of AB₅ materials. Furthermore, a description of the chill-block melt spinning process and the gas atomisation process is given.</p><p>In Part II of the thesis a digital photocalorimetric technique has been developed and applied for obtaining in situ temperature measurements during chill-block melt spinning of a Mm(NiCoMnAl)5 hydride forming alloy (Mm = Mischmetal of rare earths). Compared with conventional colour transmission temperature measurements, this technique offers a special advantage in terms of a high temperature resolutional and positional accuracy, which under the prevailing experimental conditions were found to be +/- 29 K and +/- 0.1 mm, respectively. Moreover, it is shown that the cooling rate in solid state is approximately 2.5 times higher than that observed during solidification, indicating that the solid ribbon stayed in intimate contact with the wheel surface down to very low metal temperatures before the bond was broken. During this contact period the cooling regime shifted from near ideal in the melt puddle to near Newtonian towards the end, when the heat transfer from the solid ribbon to the wheel became the rate controlling step.</p><p>In Part III of the thesis the changes of the crystal structure and the grain structure of La_0.60Ce_0.29Pr_0.04Nd_0.07Ni_3.37Co_0.79Mn_0.25Al_0.74 with increasing cooling rate during chill-block melt spinning are described. Totally, the material was rapidly solidified at 9 different cooling rates. The grain structure, crystallographic texture and the lattice parameters were studied by means of electron microscopy and powder X-ray diffraction. Additionally, the density of the rapidly solidified materials was measured by a gas pycnometer. All these properties were found to change with increasing cooling rate. The grain size decreased continuously with increasing cooling rate and was in the range of 1-5 μm. The strength of the crystallographic texture first increased and then decreased with increasing cooling rate. Transmission electron microscopy studies revealed that the grains contained a large amount of crystallographic twins and that the solidification morphology changed from cellular to plane front at a cooling rate during solidification of approximately 6·10⁴ Ks^-1. The unit cell volume and the density followed the same pattern with increasing cooling rate and decreased within each solidification morphology, but at the cooling rate from which the morphology changed, both these parameters suddenly increased. The identical variations in the unit cell volume and the density is explained by formation of excess lattice vacancies during rapid solidification.</p><p>In Part IV of the thesis rapid solidification of the materials La_0.60Ce₀.₂₇Pr_0.04Nd_0.09Ni_4.76Sn_0.24 and LaNi_4.76Sn_0.24 at 7 different cooling rates are described. The materials were analysed by means of electron microscopy and powder X-ray diffraction. The grain structures of both alloys were found to be in the nanometer range, and the grain sizes were almost invariant with increasing cooling rate. Furthermore, the lattice parameters of these materials were almost unaffected by increasing cooling rate. However, elemental line scans showed that the tin containing materials were not chemically homogeneous after chill-block melt spinning. The tin and nickel level fluctuated in an opposite manner, and the origin of these fluctuations is possibly due to inhomogeneities in the master alloys produced prior to rapid solidification.</p><p>Part V of the thesis deals with the effect of heat treatment of the rapidly solidified materials presented in Part III and IV. The first material was heat treated at 400°C and the latter two at 1000°C and 900°C respectively. Electron microscopy investigations showed that the grain structure of the first material remained unchanged during the heat treatment while the latter two were subject to sincere grain growth. The inhomogeneities were removed during the heat treatment, and X-ray powder diffraction showed that the lattice constants were changing towards equilibrium values during the heat treatment. Furthermore, the density variations in the rapidly solidified material in Part III were removed by the heat treatment. This change and the change of the lattice parameters were probably due to annihilation of excess lattice vacancies during the heat treatment.</p><p>Finally, in Part VI of the thesis the measured variations in the lattice parameters with increasing cooling rate are compared with the electrochemical hydrogen storage properties of the materials, which has been studied in a parallel work. It is shown that the hydrogen storage capacity and the absorption pressure of the material in Part III are controlled by the unit cell volume and hence the cooling rate during solidification.</p>

<p>The present work reports on the effect of Mn-, Zr- and Sc-additions upon hot deformation properties, recrystallization properties and mechanical properties for different temper conditions of Al-Mg alloys.</p><p>It can be stated that the addition of Mn, Zr and Sc improves the recrystallization properties and the mechanical properties of Al-Mg alloys. It should be emphasised that the precipitation of the metastable cubic Al₃Zr and the stable cubic Al₃(Sc,Zr) is favourable in an aluminium-magnesium matrix due to a close similarity of the lattice structures. The Al₃(Sc,Zr)-phase is similar to the equilibrium Al₃Sc-phase and has a high thermal stability and thus the coherency with the aluminium matrix is retained to very high temperatures. The present work has demonstrated the beneficial features of the Al₃(Sc,Zr)- phase upon recrystallization and strength. This also results in an increase in the deformation resistance and a reduction in the hot ductility. In particluar, manganese reduces hot ductility.</p><p>After casting most of the Zr and Sc remained in solid solution. The Mn was partly present in large primary constituent particles and partly in solid solution. Segregations of all three elements were detected. Decomposition of solid solutions of these elements resulted in the formation of dispersoids of the type Al₃Mn (orthorombic), Al₃Zr (cubic) and Al₃(Sc,Zr) (cubic).</p><p>It was found that the flow stress increased in the presence of the dispersoids. As compared to the alloy without dispersoids, the presence of Al₆Mn and Al₃Zr or Al₃(Sc,Zr) increased the flow stress by 20-100% depending on the temperature and strain rate. The effect of the particles decreases as the Zener- Hollomon parameter increases. Extrusion experiments also confirm these results. In addition, manganese reduces the hot ductility considerably.</p><p>Furthermore, the present work has demonstrated that the recrystallization properties of Al-Mg alloys may be affected considerably by introducing Mn, Zr and Sc. The recrystallization behaviour after hot deformation may be effectively determined by the Zener drag exhibited by the dispersoids on grain boundaries. Al₆Mn showed to be least effective while Al₃(Sc,Zr) is extremely effective in retarding recrystallization.</p><p>After cold deformation, however, the recrystallization behaviour is different due to a higher amount of stored energy. In the alloy without dispersoids, recrystallization occurred by classical nucleation at microstructural heterogeneities, while particle stimulated nucleation operates in the other alloys. Recrystallization of cold rolled material resulted in an extremely finegrained microstructure. Once recrystallized, extensive grain growth occurs in alloys containing Al₆Mn and/or Al₃Zr. Contrary, alloys containing Al₆Mn and Al₃(Sc,Zr) are very stable and the fine-grained structure seems to be very stable up to 550°C. This clearly proves that Al₃(Sc,Zr) are thermally stable and efficiently pin grain boundaries up to very high temperatures.</p><p>In the last part of this thesis the mechanical properties of the investigated alloys were mechanically tested in several temper conditions. It was found that the presence of Al₆Mn and Al₃(Sc,Zr) caused an increase in the flow stress of 36 MPa in the O-temper condition, probably due to the Orowan mechanism. The effect of Al₆Mn and Al₃Zr alone or in combination was less pronounced.</p><p>Furthermore, the retained deformation microstructure after extrusion was associated with the Zener drag forces exhibited by the dispersoids and resulted in considerable strengthening. For instance, the combination of Al₆Mn and Al₃(Sc,Zr) increased the strength by approximately 100 MPa compared to the dispersoid free alloy. Again the effect of Al₆Mn and Al₃Zr is less pronounced due to the lower capacity in retarding recrystallization.</p><p>The capability of the dispersoids to retard recrystallization should be an opportunity to increase the strength of the heat-affected zone after fusion welding. This is an important aspect since strain hardened conditions are used commercially. However, it has been demonstrated that a complete utilisation of the strength increase in the base material is not achieved as long as the weld metal is the weakest part in the weldment. However, a yield strength of 160 MPa was achieved for the material containing both Al₆Mn and Al₃(Sc,Zr), while somewhat lower values were obtained for the alloys with Al₆Mn and/or Al₃Zr.</p>