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ÖgeA study of protective coatings on bearings in electric vehicles: mitigating the impact of shaft voltages and currents(Graduate School, 2025-06-30)In the past decade, advances in technology have led to significant improvements in energy storage and power transmission systems. These developments have made electric vehicles more accessible and have contributed to the reduction of environmentally harmful CO₂ emissions by promoting widespread adoption. Unlike internal combustion engines, electric vehicles employ electric motors, which can result in the emergence of shaft voltages or stray bearing currents. These parasitic currents can damage the surface of contacting material pairs under operational conditions, increasing both the coefficient of friction (COF) and wear losses. To mitigate the detrimental effects of stray currents and minimize wear, various methods are employed, including grounding, the use of insulating interlayers, lubricants with high dielectric constants, and thin-film coatings. In this context, many coatings with high wear resistance, hardness, and elastic modulus have been produced by Physical Vapor Deposition (PVD) method and they have been extensively investigated. Among them, CrN and TiN coatings have gained significant attention in machining, cutting/drilling, and forming operations due to their high hardness and resistance to abrasive wear. However, their high COF and the degradation of coating properties at elevated temperatures during wear limit their service life and application range. In order to overcome these limitations, structural enhancements have been made by designing superlattice, nanocomposite or multilayer coatings, and by incorporating transition metal oxides such as those of Mo, V, Cr, and Al into CrN and TiN matrices. Their oxides form under high frictional temperatures and contribute to reducing COF due to their lubricating properties. In this study a commercial AlCrN coating which is commonly used in industrial applications and a vanadium modified one, (Cr,Al,V)N were deposited on M35 HSS by cathodic arc deposition. The structural characterization of the coatings was performed using X-ray diffraction (XRD), while their morphology and thickness were examined by scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS). The XRD results revealed the diffraction planes corresponding to NaCl type FCC structure. During wear, it is well known that Joule heating at the contact point can reach extremely high temperatures, particularly under applied current. These elevated temperatures promote the oxidation of V, Cr, and Al, which significantly affects the wear behavior. Among transition metals, vanadium oxides are particularly important due to the formation of Magnéli phases, which possess low shear strength and exhibit solid lubricating properties. The presence of such oxide phases leads to a substantial reduction in COF. To investigate the tribological behavior of the coatings under electrical effects, a modified ball on disk tribometer was used. Tests were conducted under dry sliding conditions using a 10 mm diameter 52100 steel ball as the counterface. Each coating was tested under four current levels (0 mA, 300 mA, 800 mA, 1600 mA), two sliding speeds (5 cm/s and 10 cm/s), a normal load of 2 N, and a total sliding distance of 200 meters. In order to understant the influence of electrical current on the COF, the current was applied shortly after 20 meters of sliding or when a steady-state COF was reached. Following the tests, worn surfaces were analyzed using SEM/EDS, surface profilometry, and Raman spectroscopy. Tests were repeated for each parameter for 3 time in order to obtain accurate results. Analysis of COF versus sliding distance graphs for the AlCrN coating at 5 cm/s revealed a decrease in average COF from 0.79 ± 0.10 (0 mA) to 0.61 ± 0.02 for 300 mA, 0.58 ± 0.04 for 800 mA, and 0.55 ± 0.06 for 1600 mA. At 10 cm/s, the average COF at 0 mA was 0.71 ± 0.06—relatively lower than at 5 cm/s and further dropped to 0.61 ± 0.02 (300 mA) and 0.54 ± 0.04 (800 mA), with no significant change at 1600 mA. For the (Cr,Al,V)N coating, the average COF at 5 cm/s decreased from 0.90 ± 0.08 (0 mA) to 0.50 ± 0.03 (300 mA). However, further increase in current to 800 mA and 1600 mA showed no further significant reductions in COF resulting 0.50±0.05 at 1600mA applied. On the other hand, in the case of no current application, any significant difference between sliding speeds of 5 cm/sec and 10 cm/sec in COF was not observed. However, when 300mA current was applied, the COF dropped from 0.81± 0.07 to 0.51 ± 0.04. While there was not an important change (0.51 ± 0.03) for 800 mA comparing to 300 mA, increasing the current to 1600 mA droped the COF to 0.49 ± 0.05. Regardless of coating type, the current application caused a reduction in COF. Regarding the wear behavior of the 52100 steel counter ball, wear losses increased with higher current for AlCrN at both sliding speeds. Interestingly, at 300 mA and 800 mA, increasing the speed from 5 cm/s to 10 cm/s resulted in decreased wear on the ball, whereas at 1600 mA, the trend recorded, with wear increasing alongside speed. This suggests that the combined effect of high current and speed exacerbates wear on the counterface. For the (Cr,Al,V)N coating, no significant wear occurred on the steel ball at 0 mA and 300 mA. However, at 800 mA and 1600 mA, higher speed led to increased wear, indicating that elevated current and speed together intensify thermal effects at the contact interface, thereby increasing wear damage. The examination of the wear scar with SEM revealed both adhesive and abrasive wear marks along with the pitts that result from the arc discharges. These damage features were observed to intensify with increasing applied current. For AlCrN, wear volume increased with current at both speeds. However, for a given current, increasing the sliding speed from 5 cm/s to 10 cm/s reduced wear volume: from 20.649 ± 3.1717 × 10⁻³ mm³ to 24.44 ± 1.897 × 10⁻³ mm³ at 300 mA; from 96.0108 ± 6.07 × 10⁻³ mm³ to 77.488 ± 9.7814 × 10⁻³ mm³ at 800 mA; and from 105.9554 ± 10.007 × 10⁻³ mm³ to 42.44 × 10⁻³ mm³ at 1600 mA. For (Cr,Al,V)N, no wear loss was observed at 300 mA for 5cm/s speed while there is partial removal of the coating is observed at 10cm/s sliding speed and worn volume is measured to be 8.244±4.121 × 10⁻³ mm³. At 800 mA, wear volume observe to be unchanged 4.932 ± 2.7814 × 10⁻³ mm³ at 5cm/s as recorded to be 5.263 ± 2.7814 × 10⁻³ mm³ when the speed is increased to 10cm/s to be. At 1600 mA, wear volume peaked at 103.70305 ± 17.4647 × 10⁻³ mm³ at 5 cm/s but dropped significantly to 32.533 ± 16.5057 × 10⁻³ mm³ at 10 cm/s.
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Öge26MnB5 kalite çeliğin statik deformasyon yaşlanması davranışının incelenmesi(Lisansüstü Eğitim Enstitüsü, 2025-05-29)Otomotiv sektöründe son yıllarda performans, maliyet ve karbon emisyonu gibi gereksinimleri karşılamak adına üstün mekanik özelliklere sahip yeni nesil çelikler geliştirilmektedir. Bu çelik kalitelerinden olan 26MnB5 çeliği işlenebilirlik, kaynaklanabilirlik ve yüksek dayanım özellikleri sebebi ile otomotiv sektöründe farklı parçaların imalatında kullanılmaktadır. Bu tez çalışması kapsamında iki farklı ana grupta düşük sıcaklık ve yüksek sıcaklık aralığında ticari 26MnB5 çeliğinin statik deformasyon yaşlanması davranışı incelenmiştir. Bu amaçla nominal kalınlığı 2 mm olan ticari 26MnB5 çeliği çekme deneyi numunelerine oda sıcaklığında %4 eksenel deformasyon uygulanmıştır. Düşük sıcaklık test grubunda numuneler %4 eksenel ön deformasyon sonrası 150 ℃, 200 ℃, 225 ℃, 250 ℃, 300 ℃ ve 400 ℃ sıcaklıklarda 1, 10, 100 ve 1000 dakika olmak üzere yaşlanma koşullarına tabi tutulup sonrasında çekme deneyleri yapılmıştır. Deformasyon yaşlanması sonrası ulaşılan mekanik özellikler hesaplanırken, ön deformasyon sonrası elde edilen akma dayanımı değerinden yaşlanma sonrası ulaşılan akma dayanımı değeri arasındaki fark alınmıştır. Dayanım değerlerinin hesaplanmasında numunenin orijinal kesit alanları dikkate alınmıştır. Ayrıca yaşlanma öncesi ve sonrası mikroyapısal değişimleri incelemek amacı ile optik mikroskop, taramalı elektron mikroskobu (SEM) ve X-ışını difraksiyonu (XRD) analizleri gerçekleştirilmiştir. Yüksek sıcaklık test grubunda ise numuneler oda sıcaklığında 1 hafta bekletilmiş, 800 ℃'de 5 dakika ve 950 ℃'de 25 dakika olmak üzere belirli sıcaklık ve sürelerde tutulmuş, sonrasında çekme deneyine tabi tutulmuştur. Bu deney setinde görülen sonuçlarda 26MnB5 kalite çeliğin ön deformasyon ve yaşlanma işlemi sonunda akma dayanımının arttığı tespit edilmiştir. Yüksek sıcaklık test grubunda ön deformasyon sonrası akma dayanımı değeri farkı ve yaşlanma sonrası akma dayanımı değeri farkı en fazla 800 ℃'de 5 dakika yaşlandırılan çeliğe ait iken, düşük sıcaklık test grubunda bu değer 250 ℃'de 1000 dakika yaşlandırılan çeliğe aittir. Düşük sıcaklık test grubununun ön deformasyon sonrası elde edilen akma dayanımı değeri ve yaşlandırma işlemi sonrası elde edilen akma dayanımı değerleri tespit edilmiştir. Buradan hareketle kinetik hesaplaması için Johnson-Mehl-Avrami-Kolmogorov (JMAK) modeli kullanılmış ve çeliğin aktivasyon enerjisi 163,57 kJ/mol olarak hesaplanmıştır. Deneysel sonuçlar, uygulamada ardışık deformasyon ve tavlama işlemlerine maruz kalan 26MnB5 kalite çeliğin deformasyon yaşlanmasının etkin olduğunu göstermesi bakımında otomotiv sektörüne yönelik akademik ve endüstriyen çalışmalara katkı sağlayacak nitelik taşımaktadır.
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ÖgeDesign of boron doped (nickel manganese cobalt containing) NMC 811 cathode active materials(Graduate School, 2024-11-21)Many countries have announced that they will gradually ban the sale of internal combustion engine vehicles (ICE) between 2025 and 2050 under the zero emission policy, and only zero emission vehicles (ZEV) will be sold in the near future. Nowadays, there are many brands that produce electric vehicles (EV) and they are constantly making new investments. In general, steps are being taken to improve the batteries of electric vehicles, whose biggest problem is range. According to 2023 data, lithium ion batteries have a market volume of $55 billion, and many researchers expect a compound annual growth rate (CAGR) of around 20% by 2032. Although lithium had been utilized previously in 1980, Goodenough was the first to employ a transition metal as a cathode active material in a layered (2D) structure, LiCoO2 (LCO). Then, because LCO batteries did not function at high charging rates and had a safety issue at high temperatures, due to the close ionic diameter of the Ni2+ ion and the Li+ ion, they were replaced by LiNiO2 (LNO) chemistry. When lithium leaves the cathode during charging, the Ni2+ ion fills the lithium gaps, closing the passageway of lithium. It made diffusion difficult and caused loss of capacity in the battery. Later, LiMnO2 (LMO) replaced LCO as it was economically convenient and environmentally friendly. LMO, which preserved its structure well especially at high temperatures, experienced capacity loss as a result of long cycles. Then, alternatively LiFePO4 (LFP) was utilized as cathode active material due to its environmentally friendly behavior, and high electrochemical stability at 3.5V. Later, while the specific capacity was increased by doping Ni into the LCOs, Al is added into the structure to stabilize it. Ni0.8Co0.15Al0.05 (NCA) structure offers high electrochemical performance, however, its use is restricted in some places due to security problems. NMC contains nickel, manganese and cobalt. It has been studied extensively because it offers high energy and power densities. NMC cathodes were produced in different compositions such as NMC333, NMC532, NMC622, NMC811 to optimize the capacity and the cycle life. The studies reveal that increasing Ni content in chemistry, increases the capacity of the cell but causes several problems such as chemical instability in the structures hence weak capacity retention over cycles. Today, while NMC cathode active materials can be produced with many techniques, co-precipitation method stands forward as it enables to fabricate particles with narrow particle-size distribution, high tap density with spherical morphology. The process consists of two steps: precipitation and calcination. The morphology, the structure and size of the powders obtained from the co-precipitation method are greatly affected by precipitation (pH, mixing temperature, ions concentration, mixing duration and environment) as well as calcination Therefore, in order to fabricate 5-15 micrometer sized, spherical shaped NMC811 powders via coprecipitation researchers have realized many optimization studies in the past. NMC811 suffers from low initial coulombic efficiency and capacity retention over long cycling. Structural, morphological and chemical analyses reveal that 'cation mixing', phase transformation in cycling, microcracking and hence oxygen evolution from the structure are the main problems encountered in the use of NMC811. Cation mixing is the electrochemical transformation of the crystal structure from the layered state to the rock-salt phase during the operation of the battery. Cation mixing occurs when the low-valence metal ions (Ni2+) migrate to the Li+ ion layer and replace the Li+ ions. Due to the low difference between the ionic diameters of Ni2+ (0.69Å) and Li+ (0.76Å) ions diameters among Ni2+, Co2+ and Mn2+ ions, the probability of Li+ ions being in the cation mixing with Ni2+ is higher than others. Cation mixing is not only formed during the synthesis of the material but can also be formed during the use of the battery, by redox reactions. Cation mixing (Ni2+/Li+) causes the system to be unstable thermodynamically and Ni reaches Ni2+ from high valence to low valence, causing Li and O to separate from the system, resulting in the loss of these elements and ultimately performance losses in the battery. The oxygen release from the structure can lead to safety problems since organic electrolyte systems are generally used. Structural analysis shows that NMC811 follows various phase transformation in cycling: from hexagonal to monoclinic (H1-M), from monoclinic to hexagonal (M-H2) and from hexagonal to a hexagonal structure with different lattice parameters (H2-H3). Electrochemically, the transformation from H2 to H3 phase between 4.15-4.2 V causes shrinkage around 3.7% in the c-direction in the hexagonal lattice and as a result, mechanical strain in the cathode active material, this strain causes micro cracks in the structure and leads to a decrease in cycle performance. While one of the ways to eliminate these obstacles is to add Li and O to the structure or to make a surface coating to stabilize the electrode/electrolyte interface, another solution is to control these transformations by doping the structure. Use of boron in doping becomes prominent as boron has high polarizing power due to its 3+ valence and small radius (0.098 nm.) and has strong and short bond length with oxygen leading to preventing oxygen release. Literature review reveals that two different strategies may be used to dop boron in the NMC chemistry. One is to doping boron in coprecipitation and the other is doping boron during calcination. The mechanism behind boron doping to the NMC structure during coprecipitation is that the (003) surface energy of the hydroxide is reduced compared to the {104} surfaces, thus providing the formation of primary particles oriented radially in this direction leading to the elongation of the crystallites in rod and needle shapes. Here in, the amount of boron doping is known to be crucial in the crystallization as B3+ cation may position in tetrahedral and octahedral sites (CN = 4 for 0.11 Å and CN = 6 for 0.27 Å) due to their small ionic radius. While the B3+ ion positioned in the octahedral sites will cause the cell to shrinkage in the c direction due to its small ionic radius compared to the TM ions but the B3+ ion positioned in the tetrahedral sites in the Li layer will cause expansion in the a and c direction. Moreover, the boron doping in calcination with lithium hydroxide reduces the lithium ions in the structure and prevents the development in the (003) direction, thus preventing the formation of the desired layered structure. The chemistry and amount of boron doping are quite effective in the performance of NMC 811. In NMC structures, by-products such as Li3BO3, LiBO2, Li2B4O7 are obtained as a result of the heat treatment of boron source with LiOH. It is observed that the formation of these by-products increases the cation mixing in the structure as the lithium consumption. Li-rich heat treatments are preferred in boron doping studies. In this study, a investigation is realized to investigate the effect of boron doping on NMC811 electrochemical performance. By adding boron in the coprecipitation and calcination steps of the cathode active material production process, the hypothesis put forth here is to examine the impact of the different characteristics of the B-doped NMC811 material on the electrochemical performance. No doping applied precursor named as 'NMC811OH' and boron-doped NMC811 named as 'NMC811OH1B' were successfully synthesized by co-precipitation method. NMC811OH and NMC811OH1B secondary particles with a size of approximately 10-15 μm and a spherical structure were produced. Galvanostatic tests reveal that NMC811 cathode active material without boron doping and calcined in air atmosphere (named as NMC811OH-air) delivers 160 mAh/g first charge capacity at C/10 and after 100 cycles at C/3 and C capacity retention are found to be 78% and 88% respectively, NMC811 cathode active material without boron doping and calcined in oxygen atmosphere so called NMC811OH-Ox delivers 203 mAh/g first charge capacity at C/10 and after 100 cycles at C/3 and C capacity retention are found to be 96.4% and 94.7% respectively. NMC811 cathode active material with boron doping (H3BO3) during co-precipitation and calcined in oxygen atmosphere named as NMC811OH1B-Ox delivers 188 mAh/g first charge capacity at C/10 and after 100 cycles at C/3 and C capacity retention are found to be 88% and 92% respectively and NMC811OH is mixed with LiOH and boron source (H3BO3) during calcination so called NMC811OH-Ox1B delivers 156 mAh/g first charge capacity at C/10 and after 100 cycles at C/3 and C capacity retention are found to be 88% and 89.67% respectively. I(003)/(104) ratio (inverse relationship with cation mixing) are found to be 1.22, 1.24, 1.22 and 1.17 for NMC811OH-air, named as NMC811OH-Ox, NMC811OH1B-Ox and NMC811OH-Ox1B. In all conditions calcined particles are sustained their spherical particle morphology and sizes being in the range of 10-15 μm. The galvanostatic performance shows that 1% H3BO3 doping during coprecipitation was insufficient to improve capacity retention compared to no doped NMC811OH. To further study the interaction of CAM with Li, potentiostatic cyclic voltammetry test is applied. For each electrode, peaks related to H1-M, M-H2 and H2-H3 transformations are noted, as expected. A right shift was detected in the H2-H3 peak in boron doped samples, but after 4 cycles, the loss in capacity peak intensity was found to be greater than the no doped sample in an oxygen environment and was associated with the oxygen loss in the structure and the capacity loss in cycle tests. The findings of this work highlights the importance of firstly Li/TM amount for boron source mixed with lithium hydroxide in calcination causing lithium deficiency caused by by-products or coating on the surface leading to cation mixing, poor charge-discharge capacity, capacity retention and high impedance. Boron doped via co-precipitation NMC811 sample have showed better electrochemical performance caused by better secondary particle morphology but 1% H3BO3 doping during coprecipitation is insufficient to improve capacity retention compared to undoped NMC811 sample. In the calcination processes carried out in oxygen and air, higher capacity was obtained due to the denser and more oriented primary particle structure with the presence of an oxygen-rich environment and better capacity retention due to the pore distribution in the internal structure and particle orientation.
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ÖgeYarı ergimiş tuz yöntemi ile kromit konsantresinden sodyum kromat üretimi(Lisansüstü Eğitim Enstitüsü, 2023)Kromit cevheri; metal sektörü, refrakter endüstrisi ve krom bileşikleri üretimine krom elementi sağlayan en önemli cevherlerden biridir. Doğada kübik spinal bir yapıda bulunan kromit cevheri içinde bulundurduğu oksit türü ve miktarına bağlı olarak 1500-1900 °C aralığında bir ergime noktasına sahiptir. Kromit minerali krom oksit yanında yüksek miktarda demir oksit ve görece daha düşük miktarda alüminyum silisyum ve magnezyum gibi elementlerin de oksitlerini barındıran bir katı çözelti olarak tanımlanabilir. Kromlu alaşımların ve bileşiklerin cevherlerinden üretiminde oluşan en temel ara bileşik ise kromattır. Kromat bileşikleri kromit konsantresinin alkali bir ortamda oksitlenmesi ile elde edilir. Kromat bileşikleri alkali ortamda kullanılan elemente bağlı olarak sodyum kromat veya potasyum kromat olarak üretilmesi mümkündür. Endüstride sarı kül suyu olarak da bilinen sodyum kromat kanserojen ve toksik bir bileşiktir. Ancak buna rağmen krom ve krom bileşikleri üretimi için endüstride vaz geçilemez olup; krom kaplama banyolarında, pigment olarak seramik ve cam sanayinde, ahşap oymacılık sektöründe ve deri sanayinde kullanım alanına sahiptir. Geleneksel olarak kromit cevherinden sodyum kromat üretimi döner fırında 950 -1200°C aralığında fırının içine oksijen üflenerek cevherin soda ile kavrulmasıyla oluşur. Ancak kavurma sırasında fırının içinde sodyum karbonat ile oluşan sodyum kromat ötektik bir bileşik oluşturarak reaksiyon sıcaklığının altında erir. Oluşan sıvı faz, katı fazın etrafını sararak katı gaz etkileşimine engel olup, reaksiyon verimini %80 -%85'civarına düşürür. Yöntemin düşük veriminden dolayı ortaya çıkan kromatlı atıklar ise hem çevreye hem insan sağlığına büyük risk oluşturur. Bu tehlikeyi önlemek adına reaksiyonun atıkları stokiyometrik olarak fırına tekrar beslenir. Ancak bu durum reaksiyon verimini yeterli ölçüde arttırmadığı gibi fırın hacminin de verimsiz kullanılmasına neden olur. Geleneksel kromat üretim yöntemi olan soda ile alkali kavurma metodunun verimsizliği, yüksek sıcaklık gerektirmesi ve fırın hacminin efektif olarak kullanılamaması yöntemin ekonomik açıdan olan dezavantajlarıdır. Geleneksel yöntemin dezavantajları araştırmacıları alternatif yöntemler geliştirmeye itmiştir. Bu yöntemlerden biri yarı eriyik tuz ortamında kromit cevherinden sodyum kromat eldesidir. Yarı eriyik tuz ortamının hem hidrometalurjik yöntemlere hem de pirometalurjik yöntemlere kıyasla birtakım avantajları vardır. Periyodik tabloda 1A grubu metallerinin hidroksitleri ısıtıldığında erirken önce kendi kristal sularında çözünür. Bu sayede sulu bir sistemin ulaşabileceğinden daha yüksek derişimlerde ve daha yüksek sıcaklıklarda çalışılmasına imkân verir. Ayrıca daha düşük bir viskoziteye sahip olarak difüzyonu kolaylaştırır. Ayrıca çevreye zarar vermeden kromat üretimini sağladığı için de literatürde yeşil yöntem olarak da adlandırılır. Yarı ergimiş tuz ortamında yapılan çalışmaların geleneksel yönteme kıyasla 500 -700°C daha düşük sıcaklıkta olması ve % 99'a varan yüksek verimiyle öne çıkmaktadır. Ayrıca proses kontrolünün kolaylığı ve tuz sisteminin yüksek miktarda aktif oksijen çözebilme kabiliyeti de yöntemin diğer avantajlarındandır. Bu tez çalışmasının amacı yarı ergimiş tuz ortamında kromit cevherinden sodyum kromat elde etmek olup, geleneksel alkali kavurmadan daha düşük sıcaklıkta daha yüksek verimle kromit elde etmek ve bu sayede de kromat üretiminden kaynaklı çevre kirliğinin azaltılması hedeflenmiştir. Bu hedefle tez kapsamında çalışılan kromat üretimi metodunun tuz bileşimi, süre, sıcaklık ve oksidasyon yöntemi gibi parametrelerinin taranması için farklı deney düzenekleri hazırlanmıştır. Tuz bileşimi olarak NaNO3-NaOH seçilmiş olup 325-475°C arasında değişen sıcaklığın etkisi ve reaksiyonun tamamlanma süresi çalışılmıştır. Ayrıca elektrokimyasal oksidasyonla kromit konsantresini kromata yükseltgemek için anodik akım altında çalışmayı mümkün kılan bir deney düzeneği de tasarlanmıştır. Bu kapsamda da elektrokimyasal sistemin koşullarını iyileştirmek için farklı akım yoğunlukları da test edilmiştir. Yapılan deneyler sonucunda kimyasal oksidasyonla ağırlıkça %40 NaNO3 içeren NaNO3-NaOH tuz bileşimde 425℃'de 4 saatlik reaksiyon sonucunda %98 verimle kromit konsantresi kromata yükseltgenmiştir. Reaksiyon atığında ise çözülebilir +6 değerlikli kromat bileşiklerine rastlanmamıştır. Bu sayede kromit cevherinden kromat üretimi geleneksel alkali kavurma yöntemine kıyasla daha yüksek verimde neredeyse 700°C daha düşük sıcaklıkta gerçekleştirilmiştir. Elektrokimyasal oksidasyon ile kromit cevherinin kromata oksitlenmesi için yapılan deneylerde ise %40 NaNO3 içeren NaNO3-NaOH tuz bileşimde 375℃'de 2 saat boyunca 1100 A/m2 akım uygulanan reaksiyon sonucunda %99 verimle kromat dönüşümü elde edilmiştir. Bu sayede literatürde yapılan çalışmalara kıyasla daha kısa sürede ve daha az tuz/konsantre oranında çalışılmış, elektrokimyasal oksidasyon sayesinde benzer verimde kromat dönüşümü sağlanmıştır.
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ÖgeIn-situ mechanical testing and digital image correlation of super duplex stainless steels to understand hydrogen embrittlement(Graduate School, 2022)Super duplex stainless steel contains ferrite and austenite phases in its microstructure in equal amounts, providing superior mechanical and corrosion properties to most single-phase stainless steels. The microstructure has a continuous ferritic matrix with austenite islands embedded in the ferrite phase. Super duplex stainless steel is a workhorse material used in many industrial applications where the conditions are extreme in terms of corrosion and mechanical loading. Exposure to such extremes is often associated with the liberation of atomic hydrogen. The microstructure can readily absorb formed hydrogen. However, the presence of hydrogen in high-strength microstructures is detrimental; it causes a decrease in load-bearing capacity and ductility, a phenomenon called hydrogen embrittlement. There are failure reports of super duplex stainless steel used in subsea applications with cathodic protection against corrosion. However, analysis of these failed components has shown that only duplex microstructures with an austenite spacing larger than 50 μm fractured due to hydrogen. Critical components with finer microstructures show outstanding endurance, and no failure of such duplex stainless steel has been reported. It has remained unclear why finely-grain duplex microstructures show such an exceptional resistance to hydrogen-induced cracking and why duplex stainless steel with coarse microstructure shows high susceptibility to hydrogen embrittlement. This thesis aims to understand the hydrogen embrittlement of super duplex stainless steel with a small (10 μm) and large (30 μm) austenite spacing microstructure. An in-situ mechanical testing method was developed to study the effect of hydrogen absorption and mechanical strain on the susceptibility to hydrogen embrittlement. The testing method comprises a miniature-sized tensile specimen mounted on a micro-tensile tester, an electrochemical cell for in-situ hydrogen charging, and an optical microscope with an extended focal depth. The sample was continuously slowly strained (0.005 mm/min = 4.17ꞏ10-6 s-1) while the microstructure was imaged until fracture. The specimens were either electrochemically pre-hydrogen charged for up to 72 days and then tested or tested with simultaneous hydrogen charging using self-made electrochemical cells. The results were stress-strain curves and thousands of micrographs which all provide information about the deformation characteristics of materials. Then, these images were processed with digital image correlation software and strain maps were generated to understand local strain behavior. The results have shown that hydrogen absorption caused mechanical softening in the austenite phase, while hardening was observed in the ferrite phase. In addition, the finely-grained duplex microstructure, which has more resistance to hydrogen embrittlement, developed far fewer strain heterogeneities than the coarse one. The austenite grains in the coarse microstructure became more plastically than the austenitic grains in the finer microstructure. Likewise, the ferrite became less affected due to hydrogen absorption in the fine microstructure due to more hydrogen trapping at grain boundaries. It became understood that the magnitude and number of strain heterogeneities are the main reason for hydrogen embrittlement. It also became understood that as long as the austenite phase has the capacity for hydrogen absorption and mechanical straining, the entire microstructure is protected against brittle fracture. Super duplex stainless steel contains ferrite and austenite phases in its microstructure in equal amounts, providing superior mechanical and corrosion properties to most single-phase stainless steels. The microstructure has a continuous ferritic matrix with austenite islands embedded in the ferrite phase. Super duplex stainless steel is a workhorse material used in many industrial applications where the conditions are extreme in terms of corrosion and mechanical loading. Exposure to such extremes is often associated with the liberation of atomic hydrogen. The microstructure can readily absorb formed hydrogen. However, the presence of hydrogen in high-strength microstructures is detrimental; it causes a decrease in load-bearing capacity and ductility, a phenomenon called hydrogen embrittlement. There are failure reports of super duplex stainless steel used in subsea applications with cathodic protection against corrosion. However, analysis of these failed components has shown that only duplex microstructures with an austenite spacing larger than 50 μm fractured due to hydrogen. Critical components with finer microstructures show outstanding endurance, and no failure of such duplex stainless steel has been reported. It has remained unclear why finely-grain duplex microstructures show such an exceptional resistance to hydrogen-induced cracking and why duplex stainless steel with coarse microstructure shows high susceptibility to hydrogen embrittlement. This thesis aims to understand the hydrogen embrittlement of super duplex stainless steel with a small (10 μm) and large (30 μm) austenite spacing microstructure. An in-situ mechanical testing method was developed to study the effect of hydrogen absorption and mechanical strain on the susceptibility to hydrogen embrittlement. The testing method comprises a miniature-sized tensile specimen mounted on a micro-tensile tester, an electrochemical cell for in-situ hydrogen charging, and an optical microscope with an extended focal depth. The sample was continuously slowly strained (0.005 mm/min = 4.17ꞏ10-6 s-1) while the microstructure was imaged until fracture. The specimens were either electrochemically pre-hydrogen charged for up to 72 days and then tested or tested with simultaneous hydrogen charging using self-made electrochemical cells. The results were stress-strain curves and thousands of micrographs which all provide information about the deformation characteristics of materials. Then, these images were processed with digital image correlation software and strain maps were generated to understand local strain behavior. The results have shown that hydrogen absorption caused mechanical softening in the austenite phase, while hardening was observed in the ferrite phase. In addition, the finely-grained duplex microstructure, which has more resistance to hydrogen embrittlement, developed far fewer strain heterogeneities than the coarse one. The austenite grains in the coarse microstructure became more plastically than the austenitic grains in the finer microstructure. Likewise, the ferrite became less affected due to hydrogen absorption in the fine microstructure due to more hydrogen trapping at grain boundaries. It became understood that the magnitude and number of strain heterogeneities are the main reason for hydrogen embrittlement. It also became understood that as long as the austenite phase has the capacity for hydrogen absorption and mechanical straining, the entire microstructure is protected against brittle fracture.