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(1)AGH UNIVERSITY OF SCIENCE AND TECHNOLOGY Faculty of Material Science and Ceramics Department of Ceramics and Refractories. A PhD DISSERTATION. Inconel 625 – Tungsten Carbide Composite System for Laser Additive Manufacturing Jan Huebner. Supervisor Prof. dr hab. inż. Dariusz Kata. KRAKÓW, POLAND January 2020 1.

(2) “In memory of my grandfather and for my family with all the love and support that they gave me”. 2.

(3) Acknowledgements First of all I would like to dedicate the whole effort for my family. Their support and encouragement kept me working on this Dissertation. I would like to thank my supervisor – prof. dr hab. inż. Dariusz Kata for his guidance and trust. The support and advices from him were inevitable throughout my work on this project. His experience in field of material science was great help. Many thanks to dr inż. Paweł Rutkowski. His help and patience were irreplaceable during experimental work and analysis. A great thanks goes to prof. dr hab. inż. Jan Kusiński who guided me at the start of my PhD studies and gave me valuable advices in the field of TEM and metallurgy. Over the years spent on this project I met a lot of wonderful people – without them it would be impossible to finish all planned work. Paweł and Aleksandra for unlimited access to SEM and valuable advices. Dr inż. Marta Urbańska, for translation insight during early days of my work. Anna, Karol and Mateusz for always believing in me. Jakub, for friendship and fresh point of view. Małgorzata, for encouragement that was invaluable. Adrian, for many conversations and support at the final stage of work. Special thanks for Andrzej, for organization of cooperation with outside institutions. I would also like to thank everyone who was not mentioned but was involved in any of experiments. I do remember and appreciate help and support that you offered me during this journey. At last, I would also like to thank prof. Kiyotaka Matsuura, who offered me an internship at Hokkaido University in Japan just before PhD studies. Because of that opportunity, my life is now what it is, and I am very grateful for that. Finally, I would like to quote William S. Clark:. “Boys, be ambitious!”. 3.

(4) Table of contents. Acknowledgements ..................................................................................................................................... 3 Table of contents ........................................................................................................................................... 4 List of Figures.................................................................................................................................................. 6 List of Tables ................................................................................................................................................ 13 List of symbols and abbreviations .................................................................................................... 15 Part I - Introduction.................................................................................................................................. 17 1.. Background ........................................................................................................................................ 18. 1.1. Research objectives ........................................................................................................................ 19 Part II - Theoretical................................................................................................................................... 21 2.. Turbine engines work conditions ............................................................................................. 22. 3.. Metal Matrix Composites.............................................................................................................. 25. 4.. Nickel based alloys.......................................................................................................................... 28 4.1.. Alloying elements influence on material properties ................................................ 31. 4.1.1.. Chromium and Molybdenum-corrosion resistance ......................................... 35. 4.1.2.. Aluminium and Titanium – γ’ phase formers ..................................................... 37. 4.1.3.. Niobium – γ” phase former ........................................................................................ 39. 4.2.. Grain growth and element segregation in Ni – based alloys.................................. 41. 5. Additive manufacturing as a bottom – up, three dimensional shaping method of advanced materials ................................................................................................................................. 45 5.1.. Principle of laser operation and energy source differences .................................. 47. 5.2.. Critical parameters for laser AM of materials ............................................................. 50. 5.3.. Laser Additive Manufacturing methods ........................................................................ 53. 5.4.. Laser cladding as Directed Energy Deposition method........................................... 57. 6. Laser processing of Ni – based Metal Matrix Composites with carbide reinforcement ............................................................................................................................................ 64 6.1.. Inconel 625 – WC system transformation during laser cladding ........................ 70. Part III – Experimental ............................................................................................................................ 77 7.. Experimental procedure............................................................................................................... 78 7.1.. Investigation methods .......................................................................................................... 80. 7.1.1.. Phase analysis ................................................................................................................. 80 4.

(5) 8.. 9.. 7.1.2.. Microstructural analysis ............................................................................................. 81. 7.1.3.. Microhardness and Nanohardness ......................................................................... 85. 7.1.4.. Grain size distribution of WC powder ................................................................... 86. 7.1.5.. Thermal analysis ............................................................................................................ 88. Laser cladding of pure Inconel 625 .......................................................................................... 89 8.1.. Materials – Inconel 625 powder ....................................................................................... 89. 8.2.. Inconel 625 sample preparation ...................................................................................... 90. 8.3.. Deposited material analysis ............................................................................................... 94. 8.4.. Results discussion – Inconel 625 ...................................................................................... 99. Laser cladding of Inconel 625 – WC by Yb – doped fiber laser....................................100 9.1.. Inconel 625 – WC powder mixtures preparation (DWC_1 = 0,64 µm) ................102. 9.2.. Yb – doped fiber laser apparatus ....................................................................................105. 9.3.. Process optimization and material analysis ..............................................................107. 9.4.. Results discussion - Inconel 625 – WC (DWC_1 = 0,64 µm) ....................................124. 9.5.. Inconel 625 – WC powder mixtures preparation (DWC_4 = 6,13 µm) ................125. 9.6.. Results discussion - Inconel 625 – WC (DWC_4 = 6,13 µm) ....................................139. 10. Heating rate effect on microstructure of Inconel 625 – WC system ..........................141 10.1.. DTA analysis of Inconel 625 – WC powder mixtures .........................................142. 10.2.. Microstructural investigation of DTA melted materials ...................................145. 10.3.. Hardness analysis of Inconel 625 – WC system obtained by DTA ................152. 10.4.. Results discussion - Inconel 625 – WC obtained by DTA .................................156. 11. Laser cladding of Inconel 625 – WC by CO2 laser and disc laser.................................158 11.1.. Sample preparation .........................................................................................................161. 11.2.. Microstructural investigation of Inconel 625 – WC_4 30 wt % samples ....163. 11.3.. Thermal expansion coefficient of Inconel 625 – WC system ...........................177. 12. Summary and conclusions .........................................................................................................179 Appendix ......................................................................................................................................................183 Bibliography ...............................................................................................................................................191. 5.

(6) List of Figures Fig. 2.1. Turbine blades: 1a) cooling canals inside the blade; 1b) protective coating on blade surface Fig. 3.1. Classification of composites based on dimensional orientation of reinforcement Fig. 3.2. Classification of composites based on matrix material Fig. 4.1.. Classification of Ni – based alloys based on composition [15] Fig. 4.2. Liquidus projection for 600°C, 1200°C, 1250°C of Ni – Cr – Mo ternary system with marked space representing typical amounts of these elements for commercially available alloys [25] Fig. 4.3. Phase diagram for 600°C, 1200°C, 1250°C of Ni – Cr – Mo ternary system with marked space representing typical amounts of these elements for commercially available alloys [25] Fig. 4.4. Liquidus projection and phase diagram for 750°C and 1200°C of Ni – Al – Ti ternary system with marked space representing typical amounts of these elements for commercially available alloys [25,28] Fig. 4.5. Phase diagram of Ni – Nb binary system [25] Fig. 4.6. Simplified liquidus projection of Ni – Nb – C ternary system [33] Fig 4.7. Microstructure of Inconel 625 – WC system showing different types of grain boundaries (own research) Fig. 5.1. Schematic representation of spontaneous and stimulated emission Fig. 5.2. Schematic representation of laser system Fig. 5.3. Schematic representation of SLA process Fig. 5.4.. Schematic representation of SLS/SLM process Fig. 5.5. Schematic representation of LENS apparatus Fig. 5.6. Classification of Directed Energy Deposition methods Fig. 5.7. Temperature measured during deposition of three consecutive layers [117]. 6.

(7) Fig. 5.8. Schematic representation of Marangoni effect in relation to the surface tension and temperature of melt pool: A) Surface tension gradient is negative, B) surface tension gradient is positive, C) surface tension gradient shifts between positive and negative [117] Fig. 5.9. Laser beam – material surface interaction Fig. 5.10. Correlation between G/R ratio and MPC length [117] Fig. 5.11. Columnar to equiaxial transition (CET) trend in relation to G and R [117] Fig. 5.12. Metallurgical bonding in relation of dilution A) too low dilution - no fusion to the substrate, B) optimal dilution value, C) too high dilution – keyholing [117] Fig. 6.1. ΔG of Cr, Mo, Nb and W carbides in function of temperature [134,135,144,145,136-143] Fig. 6.2. Binary phase diagram of W – C system Fig. 6.3. Non uniform distribution of coarse WC throughout the material [147] Fig. 6.4. Partial dissolution of WC in Inconel 625 laser cladded composite [152] Fig. 6.5. Partial dissolution of Cr2C3 in Inconel 625 laser cladded composite [162] Fig. 6.6. Melt pool appearance during laser cladding process of Inconel 625 [164] Fig. 6.7. Processing speed influence on porosity of the material [165] Fig. 6.8.. Vertical columnar grain growth in Inconel 625 based composites [168] Fig. 6.9. Microstructure of Inconel 625 welds after heat treatment for 8 hours in respective temperature [171] Fig. 6.10. Time – temperature – precipitation diagram for Inconel 625 [171] Fig. 7.1. General scheme of experimental procedure Fig. 7.2. PANalitycal X-Ray Diffractor Fig. 7.3. Accutom-50 (left) and Accutom-2 (right) cutting apparatus Fig. 7.4. TegraPol-21 apparatus with TegraForce-5 holder Fig. 7.5. LEICA DM2500M optical microscope Fig. 7.6. FEI Nova NanoSEM 200 Scanning electron microscope. 7.

(8) Fig. 7.7. HITACHI S-3500N (left) and FEI Inspect S50 (right) Scanning electron microscopes Fig. 7.8. JEAOL JEM-2010 ARP Transmission electron microscope Fig. 7.9. Future-Tech FM-700 hardness tester Fig. 7.10. CSM Instruments NHT50-183 nanohardness tester Fig. 7.11. Mastersizer 2000 apparatus Fig. 7.12. Grain size distribution of WC powders Fig. 7.13. NETZSCH STA 449 F3 apparatus (right); NETZSCH DIL 402C apparatus (right) Fig. 8.1. SEM image showing morphology of Inconel 625 powder from Höganäs company Fig. 8.2. RPMI 557 laser apparatus (Radomskie Centrum Innowacji) Fig. 8.3. Inconel 625 track deposition strategies Fig. 8.4. Inconel 625 samples – “pyramid” elevation on surface of SP sample Fig. 8.5. Temperature distribution during laser cladding process Fig 8.6. ST_2 sample A) good quality of Inconel 625/substrate connection, B) Substrate – Inconel 625 mixture observed Fig 8.7. F45_3 sample A) clear difference between deposited material and substrate, B) layer boundaries revealed by etching process Fig. 8.8. SEM image with EDS line analysis results of ST_2 sample Fig. 8.9. SEM image with EDS line analysis results of F45_3 sample Fig. 8.10. SEM EDS mapping of F45_3 sample, showing distribution of main alloying elements on substrate – Inconel 625 interlayer Fig. 9.1. Schematic representation of experiment stages Fig. 9.2. Morphology of Inconel 625 and WC_1 powders Fig. 9.3. Powder mixtures with 0,25% wt % resin (left) and 0,25 wt % dextrin (right) after 90 min of homogenization Fig. 9.4.. XRD results of prepared powder mixtures with 0,25 wt % of dextrin Fig. 9.5. Laser cladding head inside working chamber of JK2000FL laser apparatus 8.

(9) Fig. 9.6. Schematic representation of JK2000FL laser apparatus for deposition of Inconel 625 – WC system Fig. 9.7. Radiation pyrometer set up for temperature measurements Fig. 9.8. Temperature measurement of material surface obtained during laser cladding of Inconel 625 – WC_1 samples under 320 W laser power Fig. 9.9. Inconel 625 – 10 wt % WC samples that show effect of increased laser power and reduced distance between tracks Fig. 9.10. Inconel 625 – 10 and 20 wt % WC samples visible differences in microstructure between deposited tracks Fig. 9.11. Inconel 625 – 30 wt % WC sample – parallel to track cross-section shows different thickness of sublayers Fig. 9.12. Inconel 625 – 30 wt % WC sample: A) top of coating; B) second layer; C) coating – substrate boundary Fig. 9.13. SEM images of IW_0_64_20_320_0_8 sample – partially dissolved WC grains visible in Inconel 625 matrix Fig. 9.14. SEM images of IW_0_64_30_320_0_8 sample – fishbone – like eutectic structures formed at intergranular regions from partially dissolved WC grains Fig. 9.15. Point SEM EDS analysis of IW_0_64_20_320_0_8 sample with element composition of marked areas Fig. 9.16. Point SEM EDS analysis of IW_0_64_30_320_0_8 sample with element composition of marked areas Fig. 9.17. XRD phase analysis of IW_0_64_20_320_0_8 and IW_0_64_30_320_0_8 samples Fig. 9.18. SEM EDS mapping of IW_0_64_20_320_0_8 sample, showing segregation of elements during laser cladding Fig. 9.19. SEM EDS mapping of IW_0_64_30_320_0_8 sample, showing segregation of elements during laser cladding Fig. 9.20. TEM bright field images of IW_0_64_20_320_0_8 sample Fig. 9.21. TEM bright field image of IW_0_64_20_320_0_8 sample with diffraction pattern of 3 Ni – based matrix crystallites Fig. 9.22. TEM bright and dark field images of IW_0_64_20_320_0_8 – circular precipitates in metal matrix with EDS point element analysis for image C 9.

(10) Fig. 9.23. TEM bright and dark field images of IW_0_64_20_320_0_8 – TCP phases precipitate at grain boundary Fig. 9.24. TEM bright and dark field images of IW_0_64_30_320_0_8 – TCP phases precipitates at grain boundary, visible net of dislocations Fig. 9.25. TEM bright and dark field images of IW_0_64_30_320_0_8 – TCP phases precipitates at grain boundary Fig. 9.26. Hardness distribution thorough the coating in Inconel – WC_1 samples Fig. 9.27. Morphology of Inconel 625 and WC_4 powders Fig. 9.28. Powder mixtures with 0,25% wt % resin (left) and 0,25 wt % dextrin (right) after 90 min of homogenization Fig. 9.29. Parallel to tracks cross – section of IW_6_13_20_320_0_8 Fig. 9.30. Parallel to tracks cross – section of IW_6_13_20_320_0_8 – visible WC grains with TCP phases fishbone – like structures Fig. 9.31. Perpendicular to tracks cross – section of IW_6_13_20_320_0_8 Fig. 9.32. Representative temperature measurement on material surface for laser cladding of Inconel 625 – WC_4 samples Fig. 9.33. XRD phase analysis of IW_6_13_20_320_0_8 and sample Fig. 9.34. SEM EDS line element analysis performed on IW_6_13_10_320_0_8 sample Fig. 9.35. SEM EDS line element analysis performed on IW_6_13_20_320_0_8 sample Fig. 9.36. SEM EDS mapping of IW_6_13_20_320_0_8 sample, showing segregation of elements during laser cladding Fig. 9.37. TEM bright field images of IW_6_13_20_320_0_8 sample with EDS point analysis results Fig. 9.38. Bright and dark field TEM images with complex area diffraction pattern of 3 γ – Ni crystallites and WC grain Fig. 9.39. Hardness distribution thorough the coating in Inconel – WC_4 samples Fig. 10.1. DTA thermogram for Inconel 625 – WC_2 samples for 10°C/min and 30°C/min heating rates Fig. 10.2. DTA thermogram for Inconel 625 – WC_4 samples for 10°C/min and 30°C/min heating rates 10.

(11) Fig. 10.3. SEM images of Inconel 625 – WC_2 and WC_4 10 wt % 10°C/min sample with marked points for EDS analysis (Appendix - Table 10.2) Fig. 10.4. SEM images of Inconel 625 – WC_2 and WC_4 10 wt % 30°C/min sample with marked points for EDS analysis (Appendix - Table 10.3) Fig. 10.5. SEM images of Inconel 625 – WC_2 and WC_4 20 wt % 10°C/min sample with marked points for EDS analysis (Appendix - Table 10.4) Fig. 10.6. SEM images of Inconel 625 – WC_2 and WC_4 20 wt % 30°C/min sample with marked points for EDS analysis (Appendix - Table 10.5) Fig. 10.7. SEM images of Inconel 625 – WC_2 and WC_4 30 wt % 10°C/min sample with marked points for EDS analysis (Appendix - Table 10.6) Fig. 10.8. SEM images of Inconel 625 – WC_2 and WC_4 30 wt % 30°C/min sample with marked points for EDS analysis (Appendix - Table 10.7) Fig. 10.9. Average hardness of different phases in Inconel 625 – 10 wt % WC samples Fig. 10.10. Average hardness of different phases in Inconel 625 – 20 wt % WC samples Fig. 10.11. Average hardness of different phases in Inconel 625 – 30 wt % WC samples Fig. 11.1. Lasercell 1005 equipped with CO2 laser Trumpf TLF 3800 (Instytut Spawalnictwa Gliwice) Fig. 11.2. TruDisk 12002 laser with KUKAKR30/2 HA welding robot (Instytut Spawalnictwa Gliwice) Fig. 11.3. Lasercell 1005 equipped with CO2 laser Trumpf TLF 3800 schematic representation Fig. 11.4. TruDisk 12002 disc laser with KUKAKR30/2 HA welding robot schematic representation Fig. 11.5. Powder mixture of Inconel 625 – 30 wt % of WC_4 powder mixture with addition of 0,25 wt % of dextrin binder after homogenization Fig. 11.6. Optical microscopy images of CO2_340 sample: A, B) near coating surface area; C, D) coating interior; E, F) coating/substrate interlayer Fig. 11.7. SEM image of area inside layer in CO2_340 sample with marked points for EDS analysis (Appendix - Table 11.2) Fig. 11.8. SEM image of layer boundary in CO2_340 sample with marked points for EDS analysis (Appendix - Table 11.3). 11.

(12) Fig. 11.9. Optical microscopy images of CO2_1200 sample: A, B) near coating surface area; C, D) coating interior; E, F) coating/substrate interlayer Fig. 11.10. SEM image of near coating surface area with visible WC agglomerates in CO2_1200 sample with marked points for EDS analysis (Appendix - Table 11.4) Fig. 11.11. SEM image of area inside layer in CO2_1200 sample with marked points for EDS analysis (Appendix - Table 11.5) Fig. 11.12. SEM image of area inside layer with visible WC agglomerates in CO2_1200 sample with marked points for EDS analysis (Appendix - Table 11.6) Fig. 11.13. Optical microscopy images in DISC_320 sample: A, B) near coating surface area; C, D) coating interior; E, F) coating/substrate interlayer Fig. 11.14. SEM image of area inside layer in DISC_320 sample with marked points for EDS analysis (Appendix - Table 11.7) Fig. 11.15. SEM image of area between layers in DISC_320 sample with marked points for EDS analysis (Appendix - Table 11.8) Fig. 11.16. SEM image of large precipitate between TCP phases dendrites in DISC_320 sample with marked points for EDS analysis (Appendix - Table 11.9) Fig. 11.17. SEM image of coating – substrate boundary in DISC_320 sample with marked points for EDS analysis (Appendix - Table 11.10) Fig. 11.18. Optical microscopy images of DISC_2200 sample: A, B) near coating surface area; C, D) coating interior; E, F) coating/substrate interlayer Fig. 11.19. SEM image of coating – substrate boundary in DISC_2200 sample with marked points for EDS analysis (Appendix - Table 11.12) Fig. 11.20. SEM image typical eutectic structures inside Ni – matrix in DISC_2200 sample with marked points for EDS analysis (Appendix - Table 11.13) Fig. 11.21. SEM image of large precipitates surrounded by TCP phases dendritic structures in DISC_2200 sample with marked points for EDS analysis (Appendix - Table 11.14) Fig. 11.22. SEM image of large undissolved WC agglomerate in DISC_2200 sample with marked points for EDS analysis (Appendix - Table 11.15) Fig. 11.23. SEM image of typical eutectic TCP phases structures in DISC_2200 sample with marked points for EDS analysis (Appendix - Table 11.23). 12.

(13) List of Tables Table 4.1. Approximate atomic size ratio and solubility in about 1000°C in Nickel of various alloying elements [16] Table 4.2. Possible secondary carbides that can appear in alloys [21] Table 5.1. Absorptivity for different materials: M – metal, C – ceramic, P - polymer [84] Table 6.1. Carbides of Nb, Cr, Mo and W with their crystal structure (* unidentified) Table 6.2. Properties of Tungsten carbide Table 7.1. Measurement data for WC powders Table 8.1.. Element composition of Inconel 625 from Höganäs company Table 8.2. Fixed process parameters for Inconel 625 samples Table 8.3. Inconel 625 samples prepared by RPMI 557 laser apparatus Table 9.1. Composition of EuTroLoy 16625 alloy as declared by Castolin Eutectic Table 9.2. Compositions of powder mixtures made with WC_1 powder Table 9.3. Fixed laser cladding parameters for first batch of samples Table 9.4. Optimization of process parameters for powder mixtures with dextrin binder Table 9.5. Measured hardness of Inconel – WC_1 samples with nanohardness of two distinct phases in Inconel 625- WC_1 30 wt % sample Table 9.6. Compositions of powder mixtures made with WC_4 powder Table 9.7. Laser cladding parameters for samples with addition of WC_4 powder Table 9.8. TEM EDS point analysis of areas shown in Fig. 9.37 Table 9.9. Measured hardness of Inconel – WC_4 samples Table 10.1. List of samples subjected to DTA analysis with analysis conditions Table 10.8. Average hardness of individual phases of DTA obtained Inconel 625 -WC samples Table 11.1. Process parameters for samples prepared by use of Trumpf TLF 3800 CO2 laser and TruDisc 12002 disc laser. 13.

(14) Appendix: Table 10.2. SEM – EDS point analysis of Inconel 625 – WC_2 and WC_4 10 wt % 10°C/min sample (from Fig. 10.3) Table 10.3. SEM – EDS point analysis of Inconel 625 – WC_2 and WC_4 10 wt % 30°C/min sample (from Fig. 10.4) Table 10.4. SEM – EDS point analysis of Inconel 625 – WC_2 WC_4 20 wt % 10°C/min sample (from Fig. 10.5) Table 10.5. SEM – EDS point analysis of Inconel 625 – WC_2 and WC_4 20 wt % 30°C/min sample (from Fig. 10.6) Table 10.6. SEM – EDS point analysis of Inconel 625 – WC_2 and WC_4 30 wt % 10°C/min sample (from Fig. 10.7) Table 10.7. SEM – EDS point analysis of Inconel 625 – WC_2 and WC_4 30 wt % 10°C/min sample (from Fig. 10.8) Table 11.2. SEM – EDS point analysis of CO2_340 sample (from Fig. 11.7) Table 11.3. SEM – EDS point analysis of CO2_340 sample (from Fig. 11.8) Table 11.4. SEM – EDS point analysis of CO2_1200 sample (from Fig. 11.10) Table 11.5. SEM – EDS point analysis of CO2_1200 sample (from Fig. 11.11) Table 11.6. SEM – EDS point analysis of CO2_1200 sample (from Fig. 11.12) Table 11.7. SEM – EDS point analysis of DISC_320 sample (from Fig. 11.14) Table 11.8. SEM – EDS point analysis of DISC_320 sample (from Fig. 11.15) Table 11.9. SEM – EDS point analysis of DISC_320 sample (from Fig. 11.16) Table 11.10. SEM – EDS point analysis of DISC_320 sample (from Fig. 11.17) Table 11.12. SEM – EDS point analysis of DISC_2200 sample (from Fig. 11.19) Table 11.13. SEM – EDS point analysis of DISC_2200 sample (from Fig. 11.20) Table 11.14. SEM – EDS point analysis of DISC_2200 sample (from Fig. 11.21) Table 11.15. SEM – EDS point analysis of DISC_2200 sample (from Fig. 11.22) Table 11.16. SEM – EDS point analysis of DISC_2200 sample (from Fig. 11.23). 14.

(15) List of symbols and abbreviations 0D. Zero dimensional. 1D. One dimensional. 2D. Two dimensional. 3D. Three dimensional. AM. Additive manufacturing. ASTM. American Society for Testing and Materials. at %. Atomic percent. bct. Body centered tetragonal (structure). BPP. Beam Parameter Product. BST. Back Scattered Electrons. C/C. Carbon/Carbon Composite. CET. Columnar to Equiaxial Transition. CMC. Ceramic Matrix Composite. CMT. Cold Metal Transfer. DED. Directed energy Deposition. DLF. Directed Light Fabrication. DMD. Directed Metal Deposition. DMLS. Direct Metal Laser Sintering. DTA. Differential Thermal Analysis. DWC_X. Average grain diameter of WC_X powder. EDS. Energy Dispersive Spectroscopy. fcc. Face centered cubic. hcp. Hexagonal closely packed (structure). hex. Hexagonal simple (structure). IR. Infrared. LASER. Light Amplification by Stimulated Emission of Radiation. LBMD. Laser Based Metal Deposition. LC. Laser Cladding. 15.

(16) LENS. Laser Engineered Net Shaping. LFF. Laser Freeform Fabrication. LVD. Low Vacuum Detector. MGBs. Migration Grain Boundaries. MIG. Metal Inert Gas. MMC. Metal Matrix Composite. MPC. Melt Pool Circumference. PMC. Polymer Matrix Composite. SBT. Distance between tracks. SE. Scattered Electrons. SEM. Scanning Electron Microscopy. SLA. Stereolithography. SLM. Selective Laser Melting. SLS. Selective Laser Sintering. SGBs. Solidification Grain boundaries. SSGBs. Solidification Subgrain Boundaries. TCP. Topologically close - packed. TEM. Transmission Electron Microscopy. TIG. Tungsten Inert Gas. UV. Ultraviolet. wt %. Weight percent. XRD. X - ray diffractometry. 16.

(17) Part I - Introduction. 17.

(18) 1. Background The need for innovation in industry leads to constant evolution of materials and their production methods. Last fifteen years, showed that individual design and ability to adapt for new challenges are the major driving forces in material science. The possibilities presented by rapid development of specialized equipment offer solutions for problems that were unsolvable few years ago. The invention of additive manufacturing methods, sometimes called 3D printing – changed the world. Something that was perceived as technological curiosity, very quickly became global trend with enormous market. The new method of production for mostly polymer – based prototypes, rapidly evolved into all branches of industry. Nowadays, so called 3D printers can be easily found in both homes and huge factories. Because of additive manufacturing popularity, materials that we already know, need to be investigated for their suitability for 3D printing. Ni – based alloys are commonly used in aerospace and power industries. Because of good corrosion resistance combined with weldability and ability to be modified by addition of different phases, they can be potentially used in additive manufacturing. Lasers are excellent as energy source for metals deposition because of their higher melting point than polymers. Laser cladding is a method that uses energy delivered to material in form of laser beam in order to melt and solidify new structures, It allows for superior control of process parameters which results in very small interference in substrate together with good quality of obtained coating. Nowadays, power industry strives for innovations that could allow more efficient way of electricity production. Materials that can withstand both high temperature and chemically aggressive environments have potential to be used in turbine engines [1–3]. The easiest way to improve effectiveness of gas turbines is to increase work temperature. The blades are constantly exposed to chemical and mechanical factors which result in damaged material surface. It begins with appearance of microcavities and proceed to irreversible changes of material microstructure due to prolonged exposure to aggressive environment. Because there is no reliable and cost efficient technique of material regeneration, laser cladding of Ni – based Metal Matrix Composite was proposed and investigated in this dissertation. In order to check its potential as material for additive manufacturing,Inconel 625 – WC system was subjected to series of experiments,. 18.

(19) 1.1. Research objectives In this dissertation Metal Matrix Composite (MMC) with ceramic reinforcement in form of tungsten carbide (WC) was proposed as a material suitable for laser cladding. Usually, aggressive environment and long exposition to elevated temperatures may result in relatively short lifespan of materials used as blade turbines, boiler coatings etc. The opportunity to regenerate damaged element is economically attractive. However, conventional welding techniques such as Metal Inert Gas (MIG) or Tungsten Inert Gas (TIG), are dangerous because of large amount of heat delivered to the substrate. This leads to microstructure and phase composition changes of substrate material. In order to prevent interference within the substrate, laser cladding was proposed and investigated as potential solution. The energy delivered to material in form of laser beam causes precise heating of small area of material and then is followed by rapid cooling. Due to high processing speed, obtained microstructure is refined in comparison to conventional methods. Superior control of process parameter allows designing of material properties. Inclusion of WC particles in Ni – based metal matrix was expected to induce microstructure changes and improvement of material hardness. However, due to chemical complexity of the system, it is hard to predict its behavior during high intensity laser processing. Because of that, thorough analysis and evaluation was needed in order to properly describe how processing of Inconel 625 – WC system changes its properties. At the beginning, deposition of pure Inconel 625 coating was investigated in order to check material suitability for laser cladding. Then, experiment proceeded to next step which was introduction of WC particles as reinforcement into the Ni – based alloy matrix. The initial research objective was to obtain MMC with uniform distribution of WC in whole volume of the material. Optimization of process parameters and selection of powders were main factors that affect quality of obtained material. Fine WC (DWC1 = 0,64 µm) powder was used to prepare mixtures with different weight ratio of Inconel 625 and WC. However, due to high temperatures reached during processing, carbide particles dissolved in Ni – based matrix. This leads to next step, in which fine WC powder was replaced with other WC powder (DWC4 = 6,13 µm) in mixtures. Deposition of modified material showed, that it is possible to prevent complete dissolution of WC particles. In order to better understand composite behavior, Differential Thermal Analysis (DTA) was performed on powder mixtures containing two WC powders. Samples obtained with varying heating conditions and WC weight content were microstructurally examined by 19.

(20) means of Scanning Electron Microscopy (SEM). It allowed observation of microstructure evolution in Inconel 625 – WC system. Finally, deposition of material with high content of WC was investigated under low <500 W and high >1000 W power by using lasers with disc and CO2 laser beam sources. The examination of all prepared samples provided data that allowed to describe the material transformation during laser cladding process. Results obtained from conducted experiments answered important question: “Is Inconel 625 – WC composite system suitable for laser additive manufacturing?”. 20.

(21) Part II - Theoretical. 21.

(22) 2. Turbine engines work conditions Application of gas turbines for energy production is increasing year – by – year. Low emission of exhaust gases, relatively small sizes and ability to quick start-up are their main advantages. The rising need for clean and cheap energy made big companies like ABB, Siemens and Westinghouse invest in development of more efficient devices that can work at high temperature. Improvement from 1000 – 1150°C to 1200 – 1400°C resulted in rise of engine efficiency from 35 – 40% to 50 – 60%. It can be improved even further by rising average temperature at inlet source of heat for open turbines. For closed turbine engines there is a possibility to decrease average temperature by internal cooling channels but it is costly due to its complexity and manufacturing difficulties. Easiest solution lays in design of the material that is suitable for prolonged work at both aggressive environment and elevated temperature. The rotor blades, steering blades and discs are gas turbine parts that work at the highest temperature [4]. The requirements for the components working under such conditions includes good creep resistance and excellent tribological properties. Nowadays, two methods can be used for mass production of turbine blades: 1) Blades cooled by air flow through internal cooling channels system – because of geometrical complexity and precision needed to produce such piece, production cost is very high and often unprofitable. Internal channels are responsible for higher failure rate of parts produced by this method – Fig. 2.1a. 2) Protective coatings characterized by improved high temperature properties, at critical parts of blade. This method is efficient and cheaper in comparison to internal cooling channels – Fig. 2.1b.. Fig. 2.1. Turbine blades: 1a) cooling canals inside the blade; 1b) protective coating on blade surface 22.

(23) Metals most often used for coating of high temperature structures are Zn, Ni, Cr, Al, Sn, Cu, Ti, Mo and W because of their good corrosion resistance. Typical alloys contains: acid- and heat – resistant steels, brass, bronze, babbitt alloys and alloys: Pb – Sn – Cu, W – Co, Ni – Fe, Co – Mo, Zn – Al, Al – Si, Ni – Cr, Ni – Al, Ni – B – Si, Ni – Cr – B – Si, Ni – Cr – B – Si – C, Ni – Cr – Al – Y, Co – Mo – Cr – Si, Co – Cr – W - C and many others [5]. The combinations between metal or ceramic coating and metallic substrate material are valuable for various reasons that include improved corrosion and oxidation resistance, higher wear and erosion resistance is desirable for work in high temperature environment [6]. Materials characterized by high wear resistance are expensive because of relatively high concentration of alloying elements such as Ni, Co, Mo, V, W. Fabrication of whole parts from that alloys drastically increases production costs. Possible solutions include surface modification of selected area exposed to aggressive factors. It allows for change of material microstructure and chemical composition without interference in substrate material. Because of challenging geometry of manufactured parts, most of conventional methods of MMCs manufacturing cannot be used. Additionally, due to chemically complex composite systems, need for rapid, reliable and repeatable technique is required to ensure quality of fabricated material. Because of that, additive manufacturing (AM) method of Laser Engineered Net Shaping (LENS) also called Laser Cladding (LC) is often used to obtain superalloys coatings such as: Stellite 6, 12, 21, 156, 157, 158, Inconel 625, 718, 738, 800H [7–11]. These materials are suitable for work under high tension in temperature close to their melting points. They possess excellent corrosion resistance and are stable with increasing work temperature. The chemical composition of the nickel superalloys contains wide range of elements. Because of their role in, they can be divided into: - soluble in phase -Ni (solid solution in nickel matrix): Co, Cr, Ru, Mo, Re, W, - forming strenghtening ’ and ” phases precipitates: Al, Nb, Ti, Ta, - forming carbides: Cr, Mo, W, Nb, Ta, Ti, Hf.. 23.

(24) The technology chosen for coating deposition have huge impact on properties of wear and corrosion resistant coatings. Surface layers produced from the same material but with different techniques have varying physical and performance properties. Depending on the used methods, differences can be significant. Most common industry scale production methods of metal and ceramic coatings are thermal spraying [12,13] plasma arc welding and laser cladding [5,12,14].. 24.

(25) 3. Metal Matrix Composites Composites are a wide group of materials that consists of two or more phases. In order to be called composite, phases in material must be easily distinguished which is what differentiate them from solid solutions and mixtures. Their properties are based on individual phases as a result of the component properties (3.1) and are usually a combination of metal, polymer and ceramic materials. 𝑃𝑐𝑜𝑚 = 𝑓(𝑃1 ∙ 𝑉1 , … , 𝑃𝑛 ∙ 𝑉𝑛 ). (3.1). where: Pcom – property of composite Px – property of phase X Vx – volume content of phase X Composites usually consist of one phase described as composite matrix and one or more phases which are reinforcement. Due to different dimensional orientation we can distinguish three types of composites as shown in Fig. 3.1:. Fig. 3.1. Classification of composites based on dimensional orientation of reinforcement The 0 dimensional composites contains continuous matrix strengthened by another phase introduced as isolated particles in order to modify properties of material. This type of composite typically possess isotropic properties. In 1 dimensional composites, reinforcement appear in form of long or short fibers. This kind of composites is usually characterized by anisotropic properties which strongly depends on measurement direction. Finally, the 2 dimensional composites are built from layers of 25.

(26) different materials. Similarly to 1 dimensional composites, these materials are anisotropic. However, they usually show differences in properties in two main directions – parallel and perpendicular to layers. In some cases, short fibers in 1D composites are randomly oriented which makes them similar to 0D composites . On the contrary, when fibers are aligned in specific pattern – for example 90° between them – formation of reinforcement net results in properties similar to laminates. The other classification of composites is based on type of matrix material, as presented in Fig. 3.2.. Fig. 3.2. Classification of composites based on matrix material The type of composite matrix material mostly depends on application which it is designed for. The Polymer Matrix Composites (PMC) are suitable for low temperature environment. Because of relatively good mechanical properties combined with low density, they can be used in biomaterials, automotive and transportation industries. The Metal Matrix Composites (MMC), work conditions are more challenging – they can withstand temperatures up to 1100°C in case of superalloys. The MMCs application includes automotive, aerospace, power and chemical industries. Both Ceramic Matrix Composites (CMC) and Carbon/Carbon composites (C/C) are suitable for highest temperature environment (1500 - 2000°C). However, their application is limited due to brittleness, porosity and susceptibility to oxidation in case of non – oxide ceramics. As mentioned above, MMCs are commonly and successfully used as materials in aerospace and power industry. Production of parts for high temperature application is almost exclusively based on Ni and Co superalloys. Wide range of potential applications is guaranteed due to specific properties such as high plasticity, good corrosion and wear resistance, thermal stability and excellent weldability . In order to further improve quality and lifespan of these materials, ceramic reinforcement in form of carbides or borides can 26.

(27) be used. In MMCs, metal is responsible for shape and plasticity, while ceramics improves hardness and wear resistance of material. In following chapters, Metal Matrix Properties (MMC) on an example of Inconel 625 – WC system is discussed. Due to rapid processing of composite system and its non – equilibrium state after solidification, obtained material exhibits formation of phases that usually appear after extensive heat treatment. Introduction of W and C in form of tungsten carbide induces different behavior of alloy 625 due to enhanced element microsegregation in material.. 27.

(28) 4. Nickel based alloys Ni – based alloys usage in industry is dated back to second half of nineteenth century. However, increase of interest in these type of alloys happened at the beginning of next century. In 1902, The International Nickel Company (INCO) was formed which was followed by Haynes Stellite (now Haynes International) ten years later. Both companies were developed thanks to increased need for ballistic steels for protection armor of battleships widely used around the world during World War I. Rapid development of those companies and Ni – based alloys were positively influenced by innovations in mining industry which resulted in increased need for specialized materials at the beginning of twentieth century. Between 1920 and 1940, INCO initiated many projects that led to development of precipitation hardened alloys. Due to discovery of strength improvement of Ni alloys, the age of jet propulsion aircraft was possible. As we can see, the evolution of Ni – based materials happened during last 100 – 120 years and had crucial influence on various branches of industry [15]. Nowadays, Ni – based alloys are widely used for many specialized applications in aircraft, power, transportation, defense and petrochemical industries. Thanks to their properties such as excellent weldability, ductility and toughness at low temperatures, good corrosion resistance, high strength and stability at elevated temperatures, they fit well in different working conditions. There is no strict classification system that contain all of different Ni – based alloys. They are mostly divided by composition as shown in Fig. 4.1. [15].. Fig. 4.1. Classification of Ni – based alloys based on composition [15] 28.

(29) As it is shown, four primary groups can be distinguished [15]. Commercially pure Nickel alloys consist of over 99 wt % of Ni. These alloys have good corrosion resistance but low strength and hardness. The addition of C results in appearance of brittle carbides as a result of low solubility of C in Nickel (< 0,02%). Some of commercially pure Nickel alloys possess electrical and magnetic properties that can be beneficial in limited applications. The risk of appearance of porosity in welded material is high but can be prevented by sufficient level of shielding gas. Moreover, additions of Ti and Al limit formation of pores inside material because of promoted reaction with oxygen and nitrogen that derive from atmosphere [15]. The second group – Solid Solution Strengthened alloys, shows overall good corrosion resistance in many environments. The solubility of main alloying elements in these alloys are very high which results in their high content. This alloys exhibits different properties depending on the specific alloying element. Ni – Cu system shows complete solid solubility which provides excellent corrosion resistance for seawater. In comparison, Ni – Fe materials are used because of their electrical properties and coefficient of expansion. This provides thermal geometric stability in wide ranges of temperature. Cr, Mo and W are substitutional for Ni due to similar atomic sizes. Addition of this elements provide improved corrosion resistance and/or changes to how material behave during welding. Solid Solution Strengthened alloys shows increased strength in comparison to commercially pure Nickel alloys. However, it is necessary to chose next group in order to further improve strength of Ni – based alloys [15]. The third group consists of Precipitation Strengthened alloys which sometimes are called superalloys. As the name suggests, they contain additions of elements that induce precipitation of secondary stregthening γ’ or γ” phases within material after heat treatment. The typical elements of this kind are Ti, Al and Nb. Generally, precipitates are coherent with γ – Ni austenite matrix of alloy. Two different types of precipitates are distinguished, gamma prime γ’ – Ni3Al, Ni3Ti or Ni3(Ti,Al) and gamma double prime γ” – Ni3Nb. Properties of these alloys can be modified by optimization of heat treatment process and alloying elements content. Alloys strenghtened by γ” phase are widely used in aerospace and gas turbine applications due to low speed of formation of γ” phase in comparison to γ’ phase. The name superalloys was given because of their very high strength and excellent corrosion resistance at elevated temperature. Continuous 29.

(30) development of superalloys allowed to obtain extremally high strength and corrosion resistant materials that can be used to manufacture single crystal turbine blade [15]. The last group of so called Speciality Alloys contain materials that are precipitation and dispersion hardened. These alloys are difficult to weld due to their low ductility and high hardness. Typically they can be obtained by dispersion of fine particles of oxides, that are stable in high temperatures, throughout the material. Due to presence of oxide ceramics they possess excellent high temperature oxidation resistance. Because of that, weldability of these alloys is limited due to decrease of weld join strength. Other subgroup of Speciality Alloys is described as Nickel – Aluminides. Because of presence of NiAl and Ni3Al they posses combination of high strength and good corrosion resistance [15].. 30.

(31) 4.1. Alloying elements influence on material properties One of the most important abilities that Nickel possess in comparison to other commonly used metals, is its capability to dissolve high amount of other elements. It allows for their high possible concentration in the Ni – based solid solution. Different elements are responsible for specific material changes which strongly depends on atomic radius ratio, electron and crystal structure in relation to Nickel. If above mentioned properties are similar to Ni, then element is compatible, thus will remain in solid solution. A lot of Ni – based alloys initially designed to be single phase, have compositions that exceeds solubility limit. After exposition to high temperature for extensive amount of time, formation of brittle secondary phases occurs. These behavior is typical for alloys with high concentration of Mo such as alloy 622 or 625. Similarly, two phase alloys which contains γ + γ’ shows tendency for formation of structurally complex brittle precipitates of δ, P, µ and Laves phases during prolonged heat treatment or solidification. Appearance of these phases is both positive and negative. Generally, they exhibit properties similar to ceramic: low ductility, high hardness, brittleness and good wear resistance. This is desirable in certain applications where coating needs to be resistant to mechanical factors. However, increased amount of those phases leads to decrease of corrosion resistance and lower thermal stability which restrain high temperature applications [15]. Because of close packed face centered cubic (fcc) crystal structure of Ni matrix, diffusion rates of alloying elements are very low. This promotes high microsegregation of alloying elements during solidification. Appearance of various secondary phases can be expected when supercooling leads to solidification times between 0,01 – 10 seconds. Due to complexity of Ni – based alloys and their advanced microstructural and phase transformations, phase stability diagrams are the most effective tools available for understanding processes that occurs in these alloys [15]. If atomic size ratio of element to Nickel is close to 1, it promotes strength improvement by solid solution hardening. Table 4.1. shows approximate atomic size ratio – M/Ni and solubility in about 1000°C of various alloying elements [16]. As it is presented, Al, Ti, Mn, Nb, Mo and W are characterized by adequate size ratio of maximum about ± 0,17, in comparison to Ni. They also exhibit acceptable solubility in Ni in 1000°C. Combination of these two properties allows for solid solution strengthening. 31.

(32) If Cr or Mo are present, long term exposition to elevated temperature promotes solid state precipitation reactions. This leads to eutectic – like reactions that occurs at the end of solidification, thus formation of Cr and Mo rich secondary phases is possible. Due to low diffusivity of W and Nb, creep strength of material is improved in alloys with sufficient amount of addition of these elements. Moreover, presence of Nb induces formation of γ” – Ni3Nb precipitates. Table 4.1. Approximate atomic size ratio and solubility in about 1000°C in Nickel of various alloying elements [16] Element. Approx. atomic size ratio – M/Ni. Solubility in ≈ 1000°C [wt %]. C. 1,43. 0,2. Al. 0,85. 7. Ti. 0,83. 10. Si. 1,06. 8. V. 0,94. 20. Cr. 0,99. 40. Mn. 1,10. 20. Fe. 0,99. 100. Co. 0,99. 100. Cu. 0,97. 100. Nb. 0,85. 6. Mo. 0,91. 9. Ta. 0,85. 14. W. 0,90. 38. Additions of Ti and Al are responsible for formation of γ’ phase – Ni3Al, Ni3Ti and Ni3(Ti,Al), which are effective in improving strength of material. The γ’ – Ni3(Ti,Al) phase exhibits fcc crystal structure – the same as austenitic γ – Ni. It results in good matching (mismatch < 1%) between γ and γ’ phases. In these structures, elements like Cr, Co and Fe potentially can substitute Ni atoms [17,18]. Similarly, Nb can substitute Ti or Al in γ’ phase. Generally, γ’ phase shows tendency for ordering the microstructure during heating 32.

(33) up to 800°C which leads to improved yield strength [18,19]. Dissolution of Cr or Nb directly into the γ’ phase can lead to further solid solution strengthening [20]. The γ“ – Ni3Nb phase appears in Nb rich alloys. Nb can be substituted by Ti or Al in this phase[20]. It is characterized by body centered tetragonal (bct) crystal structure. It develops high coherency strains within the material which leads to increased strength. As a metastable phase, γ” can be easily transformed into undesirable orthorhombic δ phase, with exactly same Ni3Nb stoichiometry, during long exposure to increased temperature. The δ – Ni3Nb is incoherent with γ – Ni which causes low ductility and embrittlement of alloy. There is possibility to strengthen the material due to appearance of wide range of secondary carbides after alloy processing as shown in Table 4.2. [21]. Table 4.2. Possible secondary carbides that can appear in alloys [21] Carbide type. Lattice type. Remarks. MC. Face centered cubic. M2C. Hexagonal. M3C. Orthorhombic. M6C. Face centered cubic. M7C3. Hexagonal. Cr rich alloys, resists dissolution. M23C6. Face centered cubic. Cr rich carbide, Cr can be substituted by Fe to yield carbides with Mo or W. V rich carbide, resists dissolution, reprecipitates on secondary hardening W or Mo rich carbide of W2C type, can dissolve Cr Fe3C type carbide, M = Fe, Mn, Cr with small amount of W, Mo, V W or Mo rich carbide, may contain Cr, Co, V, wear resistant. The most common include MC, M23C6 and M6C, where M stands for metal that derives from initial alloy composition and C stands for carbon. The formation of simplest one – MC, is generally possible due to presence of W, Ta, Ti, Mo or Nb. These carbides have fcc crystal structure and appear by eutectic like reactions within γ – Ni, which is possible because of segregation of metallic elements and C to liquid. MC type carbides appear at the intergranular region due to liquid solidification. It is possible to replace MC type carbides with M23C6 and M6C during prolonged heat treatment. In alloys with high Cr content formation of Cr – rich M23C6 type carbides is observed [22]. It generally forms in temperatures between 780 – 980°C and exhibits complex cubic structure. Similarly to MC 33.

(34) type carbides, appearance of M23C6 is mainly reported as discrete carbides at intergranular region. Thanks to that, creep strength improvement is achieved by grain boundary movement restriction. Last type of carbides – M6C possess complex cubic structure and forms between 815 - 980°C. Its appearance is dependent on sufficient content of Mo and/or W in system and needs to be over 6 – 8 at. %. Formation of Topologically Close Packed (TCP) phases which consists of δ, P, µ and Laves, is possible only when highly alloyed materials are subjected to long exposure to high temperature [23]. Occasionally, δ and Laves can appear at the end of solidification due to segregation of Mo and Nb respectively [21,24]. TCP phases have complex crystal structures. If formed in solid state, they show tendency to have their close packed planes in parallel direction to the [111] γ – Ni austenite matrix. Generally, their appearance is undesirable because of decrease in strength due to depletion of alloying elements in matrix. However presence of TCP phases can lead to hardness improvement of material due to high individual hardness of these phases. Because of composition of Inconel 625 alloy, ternary Ni – Cr – Mo system, γ’ and γ” – formers alloys are discussed in following chapters.. 34.

(35) 4.1.1. Chromium and Molybdenum-corrosion resistance The Ni – Cr – Mo system is the dual purpose base for both solid solution strengthened and precipitation hardened alloys. High content of Cr guarantees excellent corrosion resistance while addition of Mo provides stabilization for many intermetallic compounds like δ and P phases which are reported to appear after welding of wide range of commercially available alloys [25,26]. Fig. 4.2. shows liquidus projection of Ni – Cr – Mo ternary system. The marked space represents typical amount of alloying elements in commercially available alloys. It is important to mention that in these alloys, solidification always begins with γ – Ni austenite matrix.. Fig. 4.2. Liquidus projection for 600°C, 1200°C, 1250°C of Ni – Cr – Mo ternary system with marked space representing typical amounts of these elements for commercially available alloys [25] Ternary phase diagrams of Ni – Cr – Mo in selected temperatures are presented in Fig. 4.3. [25]. In this system, Mo shows tendency to intensive segregation to liquid which 35.

(36) leads to appearance of intermetallic A, P, δ and σ phases at the end of solidification [24,26,27]. The formation of these phases occurs at grain boundaries. Additionally, they can also form in the solid state as a result of precipitation reactions in lower temperatures. According to presented phase diagrams, it is caused by decreased solubility of elements like Cr and Mo in γ – Ni at lower temperatures.. Fig. 4.3. Phase diagram for 600°C, 1200°C, 1250°C of Ni – Cr – Mo ternary system with marked space representing typical amounts of these elements for commercially available alloys [25]. 36.

(37) 4.1.2. Aluminium and Titanium – γ’ phase formers The Ni – Ti – Al ternary system are based on Ni – Al and Ni – Ti binary systems. Both Al and Ti promotes formation of γ – γ’ precipitation hardened microstructure in Ni – based alloys. Approximate maximum solubility of Al and Ti in Ni equals about 11% for both of these elements. It decreases with increased temperature which is the reason for precipitate strengthening reactions of secondary phases like γ’ – Ni3Al, Ni3Ti or Ni3(Ti,Al). The combined amount of Ti + Al in wide range of super alloys is below 10 wt %. This means that any further addition of one these elements will result in formation of above mentioned precipitates due to exceeded solubility limit [28]. Fcc crystal structure of Ni3Al phase shows good matching with γ – Ni which means that it can appear in wide range of compositions. Similar fcc metastable Ni3Ti can form at lower temperature which transforms into the stable form of η – Ni3Ti phase while heated. The η – Ni3Ti phase possess hexagonal closed packed (hcp) crystal structure and its formation is possible in narrow composition of about 75Ni – 25Ti at % as shown in phase diagrams in Fig. 4.4. [25]. Because it usually appears as coarse platelets, it provides only slight improvement in alloy strength. Marked area shows compositions available in most commercial alloys and indicates that solidification will begin with austenitic γ – Ni. Because of segregation to liquid, Ti and Al concentration will increase in intergranular regions. However, with high enough content of Ti and Al, solidification will end by eutectic reactions leading to formation of γ’ – Ni3Al or η – Ni3Ti. Ni – Al – Ti ternary phase diagrams shows that final composition of material mainly depends on Ti and Al concentration. The Ti – rich alloys promote formation of η – Ni3Ti while Al – rich alloys promote formation of desirable in terms of material strengthening, γ’ – Ni3Al [28]. Usually, Ni – based alloys contain high amount of Cr as main alloying element.. In temperatures below 1000°C, addition on Cr does not result in appearance of new phases. Despite this, γ’ phase is shows good solubility for both Cr, Ti and Si which exhibits desirable solid solution strengthening properties [18,29,30].. 37.

(38) Fig. 4.4. Liquidus projection and phase diagram for 750°C and 1200°C of Ni – Al – Ti ternary system with marked space representing typical amounts of these elements for commercially available alloys [25,28]. 38.

(39) 4.1.3. Niobium – γ” phase former Addition of Nb as alloying elements promotes formation of γ” – Ni3Nb phase. As shown in Fig. 4.5. the maximum solubility of Nb in γ – Ni is 18,2 wt %, the eutectic reaction occurs at Nb level of 22,5 wt %. It results in formation of γ” – Ni3Nb phase [25]. However, presence of Cr or Fe in alloy, drastically decreases maximum solubility of Nb to about 9 wt % [31]. As shown in Fig. 4.6., even low amount of C in alloy can lead to appearance of secondary NbC carbide in the intergranular region due to aggressive segregation of both elements to liquid [31,32]. Solidification in Ni – Nb – C systems starts with γ – Ni similarly to Ni – Cr – Mo [33]. In presence of Fe, Cr, Si or Mo, γ” – Ni3Nb can be replaced by intermetallic Laves phase with A2B stoichiometry, where A = Fe, Ni, Cr and B = Nb, Mo or Si. Usually, NbC forms in solid state due to long term heat treatment.. Fig. 4.5. Phase diagram of Ni – Nb binary system [25]. 39.

(40) Fig. 4.6. Simplified liquidus projection of Ni – Nb – C ternary system [33] Solidification process in Nb rich alloys generally can be simplified to three following steps (4.1-4.3): 𝐿→𝛾. (4.1). 𝐿 → 𝛾 + 𝑁𝑏𝐶. (4.2). 𝐿 → 𝛾 + 𝐿𝑎𝑣𝑒𝑠. (4.3). Formation of Laves phase is strongly promoted by presence of Nb due to L → γ + NbC reaction which occurs before L → γ + Laves during solidification. However, large amount of C addition can prevent formation of Laves phase as a result of its prior reaction with Nb in liquid. It allows to conclude that formation Laves phase can be prevented with sufficient C addition to alloy. Heat treatment of these kind of materials leads to extensive segregation and uncontrolled formation of both NbC and Laves phase [21,32]. Appearance of Laves phase can be promoted by laser processing of Inconel 625 – WC system which will be presented in Experimental part of these thesis.. 40.

(41) 4.2. Grain growth and element segregation in Ni – based alloys Solidification of solid solution strengthened Ni – based alloys results in fully austenitic material. The obtained microstructure is stable at room temperature with very low segregation that leads to local differences in element concentration at subgrain level. This can cause formation of various secondary phases at the end of solidification. The microstructure evolution during welding, which includes laser processing, shows differences throughout the material. Because of austenitic structure of Ni – based alloys, it is fairly easy to obtain good quality cross – sections after proper polishing and etching. There are three major boundary types distinguished in Ni alloys as seen in Fig. 4.7. [34,35]: a) Solidification Subgrain Boundaries (SSGBs) –shows the finest structure possible to observe with light optical microscope. They are observed as grain boundaries of cells or dendrites due to different composition in comparison to alloys matrix. The so called Scheil partitioning, describes solute redistribution at SSGBs which is responsible for differences in compositions. SSGB are crystallographically described as “low angle” boundaries with less than 5° misorientation as a result of grain growth along easy growth directions which is <100>. in case of fcc and bcc metals [34]. b) Solidification Grain Boundaries (SGBs) – shows the intersection of groups or packets. They are located at boundaries of competitive grain growth which can be easily observed because of high angle of misorientation between groups of grains. Because of that, significant amount of dislocations are located along the SGB. The solute redistribution at SGBs causes high content of impurities and precipitates to appear. Usually, weld solidification cracking of Ni – based alloys, occurs along SGBs [34]. c) Migration Grain Boundaries (MGBs) – it is the Solidification Grain Boundary that forms at the end of solidification. Due to migration of crystallographic component from compositional one. MGBs usually appears as a result of continuous welding with subsequent passes in the same area of material. It occurs in order to lower boundary energy. It is considered as high angle boundary that can cut through the subgrains located in distance of few microns from SGBs [34].. 41.

(42) Fig 4.7. Microstructure of Inconel 625 – WC system showing different types of grain boundaries (own research) Behavior of solute alloying elements during solidification can be most accurately described by Brody – Flemings equation (4.4): 𝑓. 𝑠 𝐶𝑠 = 𝑘𝐶0 [1 − 1+𝛼𝑘 ]. 𝑘−1. (4.4). Where: Cs – solid composition at solid/liquid interface C0 – nominal alloy composition Fs – fraction solid k – equilibrium distribution coefficient α – dimensionless parameter defined by equation (4.5): 𝛼=. 𝐷𝑠 𝑡𝑓 𝐿2. (4.5). 42.

(43) where: Ds – diffusivity of solute in solid tf – solidification time L – half of the dendrite arm spacing Segregation in Ni – based alloys can be described by Scheil equation (4.6) if it is assumed that there is equilibrium at the solid/liquid interface, negligible diffusion in solid and dendrite tip undercooling.: 𝐶𝑠 = 𝑘𝐶0 [1 − 𝑓𝑠 ]𝑘−1. (4.6). Many authors, that analyzed the solidification process of Inconel and Stellite alloys, have shown that the tendency of individual elements to segregate into dendrite axes and interdendritic spaces depends on their k parameter (partition coefficient) [21,26,41– 46,31,32,34,36–40]. It is described as (4.7): 𝑘 = 𝐶𝑠 /𝐶1. (4.7). Where: k – partition coefficient Cs – solid composition C1 – liquid composition The k value describes how alloying element behaves during solidification. If k < 1, element tend to segregate to the liquid. It causes appearance of that element precipitates at grain boundaries at the end of solidification. While k > 1 element shows tendency to migrate to γ – Ni matrix. It can result in element enrichment in the matrix at the expense of element depletion in intergranular region. For the Inconel alloys, the elements with similar atomic radius to Ni such as Fe, Cr, W and Co have their k values close to 1. For Mo the k value is about 0.80 ÷ 0.85 and decreases with higher Fe content in the alloy. It is even lower for Nb with k value of about 0.5 which is similarly affected by presence of Fe in the alloy. The phases with elements that segregates to liquid form at the intergranular region at the end of solidification. As a 43.

(44) result of rapid processing they can form various secondary carbides or intermetallic phases. It is confirmed that solid state diffusion of alloying elements in γ – Ni based systems during welding is negligible. Therefore, similar behavior should be expected in case of laser cladding due to even lower exposition times to high temperatures during processing. Because of that it is safe to assume that solute redistribution of elements should be mostly dependent on the partition coefficient k. Ni – Cr – Mo system exhibits complicated microstructure evolution because of potential formation of TCP phases such as σ, µ and P. This is possible due to high content of Mo which stabilizes these phases. Appearance of TCP phases is mostly undesirable due to depletion of Cr and Mo in γ – Ni matrix. It happens because of aggressive segregation of Cr and Mo to liquid. TCP formation occurs due to eutectic reactions in Cr and Mo enriched liquid at the end of solidification. Because of that, the TCP phases are rich in Mo (µ and P) or Cr (σ). It results in increased brittleness of material combined with lowered ductility and corrosion resistance. Ni – Cr – Mo alloys with small additions of Fe and W can be solidified without formation of TCP phases which is prevented by low content <15 wt % of Mo. However, increased additions of these elements will promote TCP phases formation in the fusion welding zone.. 44.

(45) 5. Additive manufacturing as a bottom – up, three dimensional shaping method of advanced materials Additive manufacturing (AM) other names includes 3D printing and rapid prototyping. The method is based on layer by layer fabrication of three dimensional structures that were initially designed in computer software. In the last years, industry found additive manufacturing as promising replacement for conventional methods [47– 49]. The development of 3D printing enabled significant cost reduction combined with higher processing efficiency. Most important advantages of AM over conventional production methods includes the ability to fabricate geometrically complex structures. It is achieved due to additive nature of the process in opposite to subtractive nature of other methods. Additionally, it prevents excessive losses of raw material which translates to better eco – friendliness due to reduced emission of waste and greenhouse gases [50–53] Despite AM development and its potential as alternative technology, there are still great need for research focused on finding new materials that are suitable for 3D printing. Complete AM process consists of many stages which includes: 3D modelling, digital data processing, 3D object construction and post processing. Overall properties of fabricated material are still mostly dependent on 3D object construction. This shows that no matter how well designed process was, the quality of obtained material is adequate to selected processing method. In order to improve material quality it is important to understand processes that occur during deposition. Thus, properly planned and executed research need to be carried out for each different material and potential application. As mentioned above AM methods are based on layer by layer deposition of materials. This could be done by: material extrusion, powder bed fusion, photo – polymerization or directed energy deposition [54]. The polymer processing by extrusion, is most commonly used in simple 3D printers. Material in form of filament goes through the nozzle which enables its melting due to heat conduction. Polymer is then cooled and solidified in air as the deposition head moves forward. Because of nonefficient heat transfer by limited contact area between material and nozzle, extrusion of molten polymer is slow. As a result, manufacturing output is low when compared to other. 45.

(46) methods. Additionally, extrusion needs constant maintenance in order to keep good quality and high resolution of fabricated elements [54]. Deposition related problems can be solved by different energy – material interaction. The material can be preplaced before processing and then followed by direct energy transfer to the desired location. Laser beam is very efficient energy source due to its high intensity. It allows for direct energy absorption by material without additional transfer medium which is often responsible for energy loss. The laser beam absorption allows for photochemical or photothermal reaction which instantly cures the material [55–57]. Because of spatially coherent light generated by laser sources, it can travel long distances without beam divergence or power loss. Additionally, optic system in laser head provides ability to focus beam spot to very small sizes which is irreplaceable advantage in AM. It allows for high precision and resolution of processing [58]. Above mentioned advantages of AM were the reason for its rapid growth over the past 15 years. The ability for relative ease in fabrication of geometrically complex structures was the main difference in comparison to conventional production methods. Global market of AM grew explosively from 0,25 billion USD up to about 9 billion USD in 2018. It is expected to reach almost 20 billion USD at the end of 2020. The market analysis exhibits strong year – by – year growth of laser shares in AM industry. It is driven by increased need for metal AM which is mostly using Yb – fiber lasers in Selective Laser Melting (SLS) and Direct Metal Laser Sintering (DMLS) methods. It single handedly shows, how important laser metal processing is in the context of technology of present and future technological advancement.. 46.

(47) 5.1. Principle of laser operation and energy source differences The term LASER is the acronym of “light amplification by stimulated emission of radiation”. Laser action begins with absorption of energy form light (photons) or heat (phonons). As a result, electron gets excited from a lower to higher energy level. It may then decay to lower energy level with (stimulated emission) or without external influence (spontaneous emission) – Fig. 5.1.. Fig. 5.1. Schematic representation of spontaneous and stimulated emission Principle of laser operation is based on stimulated emission phenomenon which is necessary for creation of laser beam. When excited electron is hit by photon P1, it decays to lower energy level with simultaneous emission of photon P2 with the same energy. As a result two photons with same energy are obtained which causes electromagnetic radiation enhancement. A material that contains many atoms in excited state may emit radiation in narrow spectrum centered around specific wavelength. In order to achieve high enough energy, laser pumping process is necessary. It proceeds by multiple reflections of laser beam by set of mirrors in order to amplify radiation by passing through gain medium. Finally, it is released through partially transparent mirror – output coupler in form of a beam. The schematic representation of this device is presented in Fig. 5.2.. 47.

(48) Fig. 5.2. Schematic representation of laser system The commercially used lasers can be classified by energy source. The most commonly used in AM are:  CO2 gas lasers  Nd:YAG solid state lasers  Yb – doped fiber lasers The CO2 lasers are one of the earliest gas lasers, invented in 1964 [59]. It typically generates wavelength of about 10600 nm which is suitable for AM of metals. It is characterized by high maximum output power (0,1 – 20 kW) which made them as excellent choice for drilling, cutting, welding and surface modification [60,61]. Because of high reliability, compactness and overall simplicity, CO2 lasers are the primary choice for precision processing. However, stability issues caused by heat generation in high power operation, CO2 devices need to be under regular maintenance [62–65]. The neodymium doped yttrium aluminum garnet laser - Nd3+:Y3Al5O12 (Nd:YAG) solid state lasers that uses rod shaped Nd:YAG crystals as a medium [66]. Similarly to CO2 they are very popular in industrial applications because of even higher compactness of devices and better beam delivery efficiency [65]. Nd:YAG lasers can work in both pulse and continuous modes while producing beam wavelength of about 1064 nm. Their maximum power output reaches up to 16 kW. Energy source used in Nd:YAG lasers can be Xenon flash lamps or diode lasers. The latter are characterized with higher power efficiency in comparison to lamp – pumped. This type of lasers are commonly used for scientific research [67–70] and optimization of processing [70–72]. 48.