Hydrogen transport membranes have gained interest from industry as an alternative to pressure swing adsorption or cryogenic distillation, and are aimed to reduce cost, equipment size, energy consumption and waste generation. Doped strontium cerate stabilized with zirconium exhibit high proton conductivity and chemical stability and is therefore considered to be a promising material for hydrogen transport membranes.The purpose of the work was to investigate if phase pure strontium cerates, stabilized with zirconium and doped with thulium, ytterbium or yttrium, could be synthesized by the Pechini method, and if the membranes exhibited sufficient density. Determination of optimal calcination and sintering parameters was also important parts in the study.SrCe0.75Zr0.20M0.05O3-d, (M=Tm,Y,Yb) were synthesized via the Pechini method, followed by calcination, pressing of green bodies and sintering of dense membranes. Phase purity of powders and sintered membranes was examined by X-ray diffraction. Surface investigation and microstructure was investigated in a scanning electron microscope. Sintering behavior and thermal expansion coefficients was determined by dilatometry. Phase pure and dense orthorhombic perovskite structured SrCe0.75Zr0.20M0.05O3-d;, (M=Tm,Y,Yb) membranes, were obtained by powder calcination at 1000 C, followed by milling and conventional sintering at 1500-1600 C. SrCe0.75Zr0.20Tm0.05O3-d demonstrated the highest density of 98,6%, when sintered at 1500 C. Sintered membranes had a average grain size in the range from 3,2 – 4,9 um. Ball milling is concluded to be of vital importance to obtain sufficient density in the membranes.The membrane characterization is limited to the methods and techniques described above. Hydrogen flux across the membrane, total conductivity, stability in reducing atmosphere and thorough investigation of thermal properties are recommended for further work.

This paper deals with alternating current (AC) corrosion of low alloy carbon steel in 3.5 wt% NaCl electrolyte. Accelerated corrosion rates have been reported when exposed to AC and the corrosion mechanism is not well understood. Electrical heating of subsea pipelines, applied to avoid hydrate formation and waxing of multiphase hydrocarbon well streams has made this topic increasingly relevant in recent years. To study the effect of AC on corrosion rates, weight loss experiments under a wide range of experimental conditions were performed. Results show that AC strongly influence corrosion kinetics of the system studied. Once AC is applied, a drop in corrosion potential and the formation of a passive iron oxide surface layer is observed. This layer is porous and believed to be the result of a rapid surface alkalization. Pitting corrosion is observed for all applied AC densities at open circuit. Weight loss- and LPR measurements conform well and the use of this technique in the presence of AC is on this basis considered valid. Corrosion rates increase as applied level of AC is increased and an approximately linear relationship has been established between iAC and icorr.

This report deals with the effect of alternating current on cathodically protected steel. AC corrosion has become relevant in the offshore industry due to the introduction of the direct electric heating system (DEH). The principle with DEH is to prevent wax solidification inside pipelines by heating them up with alternating current. This can give rise to AC corrosion.DC current densities, AC current densities and DC potential have been measured for steel samples under cathodic protection with applied AC voltage. Weight loss measurements were done of both steel and sacrificial anode samples after the experiments. There have also been focused on the corrosion kinetics of AlZnIn by plotting polarization curves and potential measurements over time with applied AC current.The weight loss measurements have shown that carbon steel is protected against AC corrosion for AC current densities lower than 1300 A/m^{2}. All steel samples had corrosion rate lower than 0,1 mm/year up to this value. SEM investigation showed that no pitting corrosion occurred at any of the steel surfaces. The corrosion rate of AlZnIn increased with increasing AC current density up to 300 - 500 A/m^{2}. At higher AC current density, the sacrificial anodes became passive which resulted in decreased protection current and corrosion rate. Polarization curves of AlZnIn have also shown large resistance when AC is introduced which is probably due to surface coverings.Study of the surface morphology of AlZnIn have shown that more localized corrosion occurs with increasing AC current. At 500 A/m^{2} pitting corrosion occurred at the whole surface and the pits became deeper with increasing AC. This resulted in decreased efficiency of the sacrificial anode and consequently reduced DC current density and corrosion rate of the sacrificial anode.Potential and current measurements have shown that the galvanic potential of steel and AlZnIn became more electropositive with increasing AC current density. It is believed that this is due to more localized corrosion with increasing AC voltage and consequently lowering of the anode/cathode ratio. This will force the potential to more anodic values. The initial DC current density increased with increasing AC voltage for all samples but the passivation of sacrificial anode resulted in high reduction rate with time. Polarization curves of AlZnIn showed that the corrosion potential decreased and the corrosion current increased when AC was introduced. Within the samples exposed to AC there were no significant difference in corrosion potential and corrosion current. The polarization curves of all samples that were exposed to AC were characterized by high ohmic resistance which was due to hydrogen evolution and formation of corrosion products at the surface.

Copper are used as an alloying element in aluminum alloys to increase the strength of the material. By mixing copper and aluminum the good corrosion resistance of the pure aluminum decreases giving the alloy a lower corrosion resistance. After years of investigation on corrosion of aluminum alloys several results have shown increasing corrosion rates of aluminum that have been alloyed with both copper and zinc, giving a “grainy appearance” on the surface of the alloy. By adding copper to the aluminum increased intergranular corrosion and preferential etching of surface grains has been found after heat treatment and etching of the alloy.Before etching the aluminum alloys in alkaline and acidic environments, Glow Discharge Mass Spectroscopy (GDMS) measurements was done for all the alloys to determine the correct amount of alloying elements and impurities in each alloy. Preliminary etching trials have been performed on pure aluminum (Al 5N), aluminum containing 10 ppm copper (AlCu10), aluminum containing 100 ppm copper (AlCu100), and aluminum containing 1000 ppm copper (AlCu1000), as well as corrosion potential tests. Further, the surface of all the etched test specimens has been investigated in light microscope, Scanning Electron Microscope (SEM), Electron Backscattered Diffraction (EBSD), Confocal Microscope (IFM), Energy Dispersive Spectroscopy (EDS),and Glow Discharge Optical Electron Spectroscopy (GD-OES). These methods have been used to determine height differences between etched neighboring grains, the crystallographic orientation of grains, surface roughness, detecting elements occurring on the surface of etched specimens, and giving a depth profile of the etched specimens.For the alkaline etching trials, an increasing etching rate was found by increasing the amount of copper in the alloys, however the AlCu10-alloy showed little or no difference from the pure aluminum both in the etching trials and the corrosion potential tests. By increasing the copper amount in the aluminum the corrosion potential increased drastically for AlCu100-alloy and the AlCu1000-alloy. Further an increasing etching rate of grains having close to [111] crystallographic orientation could be seen in alkaline environment, a difference in the surface roughness between grains with different crystallographic orientations could also be seen for the alkaline etched test specimens. The GD-OES investigations done on test specimens that had been alkaline etched revealed differences in their depth profiles dependent on the etching temperature and how the specimens was treated after being etched. AlCu1000-alloys etched in acidic environment showed a different type of etching and surface after etching than the other alloys etched in the same environment. A clear difference could also be seen between the alkaline and the acidic etched AlCu1000 specimens.

The predominant site for the initiation of pitting on austenitic stainless steel has been shown to be sulfide inclusions and notably the manganese types of sulfides. Dissolution of inclusions has been observed and suggested to be the initial step for pit initiation, though several explanations for the mechanisms causing initiation has been proposed. Regarding welded stainless steels, several microstructural changes have been described and suggested to contribute to the decreased corrosion resistance in the weld zone. An area which has not been investigated much is the contribution of inclusions, in particular MnS inclusions, to the reduced corrosion resistance and the initiation of pitting in the weld zone. In the present work a literature review of the investigations focusing on the initiation of pitting by sulfide inclusions in austenitic stainless steels has been provided. A literature review of investigations focusing on the effects of welding on the microstructure, inclusions and the corrosion properties in the weld zone of austenitic stainless steels has also been provided.Experimental work was performed to obtain results which could be compared to or verify findings and suggestions from the reviewed literature regarding the initiation of pitting by inclusions. Observations of MnS inclusions as the preferential site for pit initiation in austenitic stainless steel was tried recreated. The hypothesis saying that dissolution of MnS inclusions occur prior to the initiation of pitting was also tried verified. It was also performed experiments to obtain results which could give knowledge about the effects of MnS inclusions on the corrosion properties and the initiation of pitting in the weld zone of an austenitic stainless steel. Along with this, the corrosion behaviour and microstructural changes in the weld zone in general and compared with the base metal were studied. A part of the work was also to study the effect of a lacking inert shielding gas during welding.Samples of a 316L stainless steel were first examined in SEM to study the microstructure and to identify inclusions. Then, some samples were welded, followed by SEM examination. Then all samples were polarized electrochemically in synthetic seawater with the purpose of initiating pitting. After the polarization, all samples were examined in SEM again to correlate pitting attack to inclusions and to compare microstructural changes and the corrosion behaviour in the base metal and the weld zone.It was verified that the main initiation site for pitting in the base metal was MnS inclusions. It was indicated that dissolution of the MnS inclusions started the pit initiation process, with the contribution of released compounds from the inclusions. An average percent of inclusions showing an inactive behavior regarding the initiation of pitting was 38 %, verifying similar observations reported in earlier investigations. A nitric acid treatment to remove MnS inclusions improved the pitting resistance significantly.The welding caused the initiation of corrosion at lower potentials when performing electrochemical polarization in synthetic seawater. The corrosion mechanism was probably grain boundary corrosion caused by precipitation of chromium carbides in the grain boundaries. A lacking inert shielding gas during welding did not have any effect on the potential for the initiation of corrosion when comparing two samples welded with and without shielding gas.It was indicated that compositional changes had occurred for MnS inclusions in a certain distance from the fusion line. These changes may have caused the inclusions to be more prone to initiate pitting. Further investigation should be performed to clarify the behaviour of such inclusions, and their susceptibility towards the initiation of corrosion. Grain boundary corrosion in various grades had occurred in the HAZ in certain distances from the fusion line. In a further distance from the fusion line pitting-like corrosion in clusters, possibly induced by chromium depletion after forming chromium oxides combined with a thin oxide film, had occurred. The grain boundary corrosion and the pitting cluster attack are from the experimental results believed to be more important for the failure of the welded stainless steel investigated in this work, than pitting initiated at MnS inclusions changed by the welding process.

The main object was to investigate the effect of donors, thermal donors and defects on the lifetime on n-type Czochralsky (Cz) silicon. Cz is a technique for making monocrystals by dipping a monocrystalline seed into the melt, and pulling the ingot up as the melt solidifies. Samples were prepared by sawing, grinding and polishing. Copper decoration and preferential etching was done to reveal the defects. Resistivity measurements were performed to investigate the donor distribution. The lifetime and oxygen concentration was measured.The lifetime may be influenced by donors and traps. Traps may be metallic or non-metallic. Dopants and metal contaminants are usually Scheil distributed, which typically involves a relatively stable distribution in most part of the ingot, and a rapid increase in concentration close to the bottom. Metal contaminants may be introduced from the feedstock, and dopants (either n or p dopants) are added to the melt.Oxygen is dissolved from the crucible. The oxygen concentration typically decreases with increasing length, and is controlled by the rotation speed. This was confirmed by the interstitial oxygen measurements. The oxygen concentration showed no change after heat treatment, but the presence of oxygen in itself may contribute to produce lifetime reducing agents, such as precipitates and thermal donors. Thermal donors were found in the first 18cm of the ingot. Both the resistivity and lifetime measurements was used to calculate estimated lifetimes and donor concentrations, respectively. This estimated values were compared with the measured ones, revealing that the lifetime in the ingot was SRH dominated. Thermal donors were found to have less influence on the lifetime than the trapping. Oxygen precipitations may explain the low lifetime found in the center of the ingot up to about 30cm from the crown. Such precipitations are often gathered in defect bands in the sample. CDI of a Cu decorated sample revealed a pattern of different defect densities, which also showed up at the surface of the sample after preferential etching. An H- or L- band was believed to be situated here, which both consist of small oxygen precipitates. Defects arize as a funktion of the pulling parameters, and a low growth rate was connected to the low lifetime here. Indications of the presence of non-metallic traps were found throughout the ingot. However, oxygen precipitates could not explain the non-metallic traps that was found on the outside of the ingot, and the cause of these is therefore unknown. It is however likely that these traps may be connected to the pullrate/temperature gradient-ratio. Most likely metallic traps were also present. It is uncertain whether the non-metallic or metallic traps were the most important to the lifetime of this ingot, as the relative distribution of the two cannot be decided with certainty from the results in this thesis.Contrasts in CDI and steep curves using PCD measurements vs length plots were believed to be due to the lifetime degrading agent that changes the fastest. Metallic traps and the phosphorous dopant are known to be Scheil distributed, hence they appear transparent for the lifetime measurements in areas where the distribution is stable. Thermal donors and non-metallic traps are thus more likely to create contrast in the top parts of an ingot. However, it is not necessarily the phenomenon that causes the contrast that is the major lifetime degradation contributor.

I denne masteroppgaven ble det syntetisert iridiumbaserte katalysatorer på en bærerstruktur for oksygenutviklingsreaksjonen i en PEM vannelektrolysør, ved hjelp av en polyolmetode. Katalysatorpulverne ble karakterisert ved hjelp av både elektrokjemiske (syklisk voltammetri, polarisasjon og impedans), og fysiske metoder (SEM, TEM, BET, EDS, and XPS). I tillegg ble det konstruert en apparatur for måling av elektronisk ledningsevne, og denne ble benyttet for å sammenlikne den elektroniske ledningsevnen til to ulike bærermaterialer, antimondopet tinnoksid (ATO) og wolframoksid (WO3).Målet for arbeidet var å studere effekten som bruk av bærermaterialer og ulike synteseforhold har på de elektrokatalytiske egenskapene til iridiumbaserte katalysatorer på bærere.Resultatene fra dette arbeidet viser at av de to valgte bærermaterialene så har ATO de mest egnede egenskapene for bruk som bærermateriale i et elektrodesystem, i det at den elektroniske ledningsevnen var mye høyere og det spesifikke overflatearealet var større enn hos WO3. Karakterisering av de syntetiserte katalysatorene på ATO viser elektrokatalytiske egenskaper som er på samme nivå som, eller bedre enn, andre katalysatorer laget til samme formål. Forandringer gjort på synteseforholdene (pH og konsentrasjon av materialer) virket ikke å ha noen effekt av betydning på verken morfologi, partikkelstørrelse eller den katalytiske aktiviteten til katalysatorene.

As the silicon content in a silicomanganese alloy increase, silicon carbide becomes the stable carbon phase. Little work is published on the formation of silicon carbide in the SiMn process. This thesis examines the formation of SiC through the reaction between slag, metal and coke. The goal of the thesis has been to determine where and how SiC is formed in the silicomanganese process. Focus has been given to formation through liquid-solid reactions.The investigation was carried out by heating metal and slag together with coke. XRD analysis of SiC formed in the process, as well as wettability testing of slag and metal toward SiC and graphite substrates were also examined. The raw materials consisted of two industrial SiMn alloys, named Metal 1 and 2 in the thesis. These had 19 and 28 wt.% Si, respectively. Two slags, Slag 1 and 2, were used in combination with the metal. SiC was found to form on the coke particle through reaction with both slag and metal. The formation was fastest through the metal phase and the effect of increasing silicon content in the metal was evident. A temperature increase from 1600 C to 1650 C resulted in less SiC formed on the coke particle surface for Metal 1, but gave more SiC at the surface of the coke particle for Metal 2. The decrease in SiC on the coke particle for Metal 1 is likely to be caused by a decreased driving force, as the coexistence point of C and SiC increases with increasing temperature. For Metal 2 the relative distance is smaller, making the effect neglectable. The formation of SiC occurs at the surface of the coke particle. The metal samples indicate that carbon diffuses through the SiC layer to react with silicon in the alloy. Formation through the slag is likely to go through the reduction of SiO2 to Si, before reacting to form SiC. The XRD analysis determined beta-SiC to be the dominant phase formed through both slag and metal. alpha-SiC was found in one of the industrial slag samples, indicating temperatures higher than 1700 C. The wettability testing showed both slag and metal to be non wetting toward the SiC substrate. All angles were stable with increasing temperatures. Both slags were found to be non wetting toward graphite. The wetting angle of Slag 2 was stable with increasing temperature. Slag 1 showed decreasing wetting angle with temperature, before an abrupt change back to an angle larger than the initial angle. The sudden change is likely to be caused by a reaction product, possibly SiC, forming between the substrate and slag.

Denne oppgaven har tatt for seg et 420 MPa HSLA stål. Stålet ble sveisesimulertfor å oppnå mikrostruktur tilsvarende grovkornet og interkritisk grovkornetHAZ. Interkritisk grovkornet HAZ ble oppnådd ved en tosyklus sveisesimuleringhvor materialet ble varmet opp til 1350 °C og avkjølt med Δτ8-5 = 15s og deretter gjenoppvarmet til 780 °C og avkjølt med Δτ 6-4 = 17s. Grovkornet HAZ ble oppnådd ved ensyklus sveisesimulering hvor materialet ble varmet opp til 1350 °C og avkjølt med Δτ 8-5 = 15s.Et eksempel på EBSD teknikken er vist ved en sammenligning av sekundærelektronbilde, fasekart og Image Quality kart. Sekundærelektronbilde og EBSD skannet er fra nøyaktig det samme område. På denne måten blir det demonstrert at EBSD teknikken er et verdifullt verktøy for å skille ulike stålfaser fra hverandre.Det ble så fortatt en grunnleggende EBSD karakterisering av de to sveisesyklusene. Fire skann med et totalareal på 150 µm x 150 µm ble gjort på hver av de to prøvene. Det ble konkludert med at den største forskjellen var øyer av restaustenitt som i gjennomsnitt utgjorde 2,45 % av det skannede område på tosyklus sveisesimulert prøve. I ensyklus sveisesimulert prøve var disse øyene så godt som fraværende.For tosyklus sveisesimulert stål ble det også gjennomført in situ strekktesting med påfølgende EBSD karakterisering ved -60 °C og romtemperatur. En spesiallaget kaldfinger ble videreutvilket for å kjøle ned prøven til -60 °C, da det ikke lykkes å nå -60 °C slik den opprinnelig var konstruert. EBSD karakterisering ble foretatt uten deformasjon og ved 1 %, 2 %, 3 % og 4 % forlengelse av stålet. En sammenligning av resultatene ved -60 °C og romtemperatur ble gjort. Det ble konkludert med at øyer av restaustenitt transformeres til martensitt som følge av plastisk deformasjon i større grad ved -60 °C enn ved romtemperatur som følge av høyere termodynamiske drivkrefter.Ved in situ undersøkelser må strekkbord og kaldfinger monteres i mikroskopet, dette fører til at det tar mye lengre tid å pumpe vakuum enn ved normale omstendigheter. Det tar 50 minutter å pumpe 2,5 x 10-6 mBar vakuum i tomt prøvekammer. Tilsvarende vakuum ved in situ undersøkelser tar over 5 timer å pumpe, grunnet større overflate inne i mikroskopet. Til slutt ble det gjennomført en sammenligning ved bruk av 1° og 2° Θ steglengde ved indisering av EBSD rådata. Indisering med 1° Θ steglengde ga 1,5% austenitt, mens 2° Θ steglengde ga kun 0,9% austenitt. Konklusjonen var at 2° Θ steglengde ikke oppdager de minste kornene på grunn av for grov omregning i Hough transformasjonen.

Ved sveising av HSLA-stål (High Strength Low Alloy) designet for lavtemperatur applikasjoner kan det dannes lokale sprø soner som martensitt/austenitt-faser (MA-faser). Det er derfor ønskelig å karakterisere austenitten som dannes ved tosykel sveising for å gi økt kunnskap om endringene som oppstår i mikrostrukturen.Denne masteroppgaven tar for seg karakterisering av et 420 MPa HSLA-stål ved in situ varmebehandling opp til tofaseområdet for austenitt og ferritt i kombinasjon med diffraksjon av tilbakespredte elektroner (Electron Backscatter Diffraction - EBSD). Målet med oppgaven var da å karakterisere endringer som oppstår i mikrostrukturen til stålet under in situ oppvarming i mikroskopet. Oppvarmingen ble utført ved å videreutvikle et tidligere varmebord. Stålprøvene ble ènsykel sveisesimulert til 1350 °C og avkjølt med Δt8/5. Den in situ varmebehandlingen tilsvarte da den andre sveisesykelen i en tosykel sveisesimulering. Ved prøvepreparering ble prøvene elektropolert med to ulike elektrolytter; Struers A2 og AC2. I oppgaven ble det benyttet offline-EBSD da det var nødvendig med høy skannehastighet slik at tiden for hvert EBSD-skann ble så liten som mulig.For å nå tofaseområdet til HSLA-stålet var det nødvendig å varme opp til en prøvetemperatur på over 740 °C. Da ovnen har en maksimaltemperatur på 800 °C ble det benyttet en vakuumkompatibel platinamaling mellom prøven og ovnen i varmebordet for å øke varmeledningen. Temperaturforskjellen mellom prøven og ovnen ble da reduseres helt ned til 10 °C. Det ble varmet opp til prøvetemperaturer på 750 °C, 760 °C, 770 °C, 780 °C og 790 °C. Det var nødvendig å vente i 10 minutter etter å ha nådd ønsket prøvetemperatur før det ble kjørt EBSD på grunn av termisk ekspansjon av prøven.Austenitten som ble dannet under oppvarming vokste frem langs korngrenser og hadde en orienteringssammenheng med BCC-matriksen rundt som oppfyller Kurdjomov-Sachs. Det ble observert at austenitten som dannes gradvis forsvant igjen ved økende holdetid ved høy temperatur som følge av avdamping av karbon fra prøveoverflaten. Ved oppvarming viste det seg at det ble dannet en mye større andel austenitt i prøver elektropolert med Struers AC2 enn den med Struers A2. Det ble da konstruert fasediagram for tofaseområdet til stålet på bakgrunn av andel dannet austenitt ved oppvarming. Disse viste at prøver elektropolert med Struers A2 gav mest korrekt andel austenitt.