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Behavior of TiN inclusions and their influence in random grain formation in Ni-based superalloys Soundararajan, Rajeswari 1998

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BEHAVIOR OF T i N INCLUSIONS A N D THEIR INFLUENCE IN R A N D O M G R A I N FORMATION IN Ni-BASED SUPERALLOYS by R A J E S W A R I S O U N D A R A R A J A N B . S c , Madras University (INDIA), 1990 M.Sc. , Anna University (INDIA), 1992 A THESIS S U B M I T T E D IN P A R T I A L F U L F I L L M E N T O F T H E R E Q U I R E M E N T S F O R T H E D E G R E E O F M A S T E R O F A P P L I E D S C I E N C E in T H E F A C U L T Y O F G R A D U A T E S T U D I E S (Department of Metals and Materials Engineering) We accept this thesis as conforming to the required standard T H E U N I V E R S I T Y O F BRITISH C O L U M B I A October 1998 © Rajeswari Soundararajan, 1998 In presenting this thesis in partial fulfilment of the requirements for an advanced degree at the University of British Columbia, I agree that the Library shall make it freely available for reference and study. I further agree that permission for extensive copying of this thesis for scholarly purposes may be granted by the head of my department or by his or her representatives. It is understood that copying or publication of this thesis for financial gain shall not be allowed without my written permission. Department of The University of British Columbia Vancouver, Canada Date DE-6 (2/88) Abstract Studies on the columnar to equiaxed transition and random grain formation in superalloy turbine blades have suggested that non-metallic inclusions may be a possible nucleation source. Though there is no strong evidence to substantiate the idea, during directional solidification if the inclusions are able to survive in the region ahead of the solidification front, they could act as potential nucleation sites and may aid in the formation of random grains. O n the. basis of the Turnbull-Vonnegut lattice dis-registry model, the lattice mismatch between the F C C Ni-matrix and a typical cubic TIN inclusion is determined to be less than 20%. Hence, a T i N particle can act as a potential heterogeneous nucleant in Ni-based superalloys with sufficient undercooling. The undercooling required in the Ni-based superalloys for T i N inclusion is calculated to be about 19 °C, which is comparable to the undercooling found in some single crystal superalloy turbine blades. Therefore, if the T i N particles are stable in the region ahead of the liquidus isotherm, they may be potential nucleation sites and hence cause the formation of random grains. This thesis reports on a quantitative study to verify the hypothesis that the T i N particles are stable in the region ahead of the solidification front during directional solidification. Thermodynamic calculations concerning the solubility of nitrogen and the formation of T i N in alloy IN718 have determined that the equilibrium nitrogen content required for the formation of T i N is about 39 ppm [N] (at Tii q = 1 3 4 0 ° C ) . Experimental studies on alloy IN718 at various conditions have indicated that T i N precipitation will not take place once the equilibrium nitrogen content is below this value of 39 ppm at Tu q . In the solid/liquid mushy region, the combined effects factors such as: a) cooling, b) segregation of T i and nitrogen, and c). rejection of nitrogen into the bulk liquid, reduce the solubility limit to below 39 ppm. Hence T i N precipitation takes place in the segregation zone. The experiments have established that T i N particles precipitated in the solid/liquid mushy zone will not float out of the interface. Therefore, the nitrogen content must ii exceed the saturation solubility of T i N at Tu q to provide nuclei for random grain formation. Due to nitrogen rejection into the bulk liquid, there is T i N precipitation just above the T i i q . Directional solidification (DS) experiments on IN718 samples under nitrogen pressure and at a withdrawal rate of 2.4u.m/sec have revealed that T i N particles of size > 10 microns (which have a rise velocity > 28 u,m/sec) are able to float and eventually get collected at the top of the ingot (which is the final portion to solidify). These might be the T i N particles that are precipitated in the liquid ahead of the solidification front, i.e. above the Tii q , and which subsequently undergo flotation. Smaller T i N particles of size < 4 microns are found at the bottom portions of the ingot which indicates that these particles may be precipitated later during solidification in the solid/liquid mushy zone, and thus didn't "escape" out of the interface. Hence, the withdrawal rate should be at least an order lesser in magnitude than the velocity of rise of the T i N particles for substantial flotation to occur. iii Table Of Contents Abstract ii Table of Contents iv List of Figures vii List of Tables xii List of Symbols xiii Acknowledgements xv Introduction 1 Chapter 2: Literature Review 4 2.1 Random grain formation 4 2.1.1 Fundamentals o f C E T 5 2.1.1.1 Dendrite Undercooling 7 2.2 Spurious Grain Formation in Ni-Based Alloys 9 2.2.1 Thermosolutal convection 9 2.2.2 Constitutional Undercooling 14 2.3 Criterion For Spurious Grain Nucleation 15 2.4 Inclusions and Random Grain Formation 20 2.5 Thermodynamics of Inclusion Formation 25 2.5.1 Nitrogen in Superalloys 25 2.5.1.1 Solubility of Nitrogen 25 2.5.2 Chemistry of T i N Precipitation 27 2.5.3 Studies in Alloy IN718 29 2.6 Summary 32 Chapter 3: - Objectives 34 Chapter 4: Theoretical Methodology for T i N Formation in Alloy IN718 3 5 4.1 Solubility of Nitrogen in Alloy IN718 35 4.2 Solubility of T i N in Alloy IN718 36 iv Chapter 5: Experimental Studies 41 5.1 Alloy Compositions used for Experiments 41 5.2 Sample Preparation 41 5.3 Apparatus and Procedure 43 5.4 Metallographic Analysis 47 5.5 Primary Dendrite Arm Spacing 51 5.6. Buoyancy Effect O n T i N 54 5.7 Experiments 57 5.7.1 Nitrogen Analysis of the D S cast Samples 60 Chapter 6: Results 61 6.1 Directional Solidification and Quenching Experiments 61 6.2 Directional Solidification experiments 63 6.2.1 Samples 90 ppm [N], and 64 ppm [N] 63 6.2.2 Homogeneous Sample Preparation 66 6.2.3 90 ppm [N] samples D S at 3.5 x 10"5 Torr 66 6.2.4 90 ppm [N] D S at 100 microns Nitrogen Pressure 68 6.2.5 90 ppm [N] Sample D S at 400 microns Nitrogen Pressure 72 6.2.6 D S Experiments With 20 ppm [N] Samples 76 Chapter 7: Discussions 80 7.1 Directional Solidification and Quenching Experiments 80 7.2 Directional Solidification (DS) Experiments 80 7.3 Summary of T i N Particle Distribution and Nitrogen Analysis 83 7.4 T i N Particle Distribution 93 Chapter 8: Conclusions and Recommendations 95 8.1 Summary 95 8.2 Recommendations For Future Work 100 References 102 v Appendix 1 109 Appendix 2 111 Appendix 3 113 Appendix 4 115 Appendix 5 119 Appendix 6 121 Appendix 7 124 Appendix 8 128 v i List of Figures Figure 1. Fatigue properties of IN718, illustrating the influence of inclusions on premature failure at T = 350 °C; strain amplitude = 0.62%. (from [1]). 3 Figure 2. The three regimes of grain defect formation and the dependence on dendrite arm spacing, (from [5]). 10 Figure 3. Isolated mis-oriented columnar grains during directional solidification of alloy SX-1 (after [5]). 11 Figure 4. Complete breakdown of single solidification and blockage of solidification Front by equiaxed grains, (after [6]). 11 Figure 5. A typical freckle chain observed during solidification of alloy SX-1. (from [6]). 12 Figure 6. The dependence of freckle chain formation on primary dendrite arm spacing, (after [6]). 13 Figure 7. Depending upon the undercooling capacity of the given alloy a macroscopic concave curvature of the liquidus isotherm can cause the formation of stray grains (a) or result in fast lateral growth of secondary dendrites (b). (from [35]). 16 Figure 8. Schematic diagram showing the growth of a secondary dendrite arm into a platform in the presence of a transverse temperature gradient. (from[6]). 17 Figure 9. Effect of transverse temperature gradient in alloy C M S X - 4 . (from [6]). 19 Figure 10. Lattice mis-match vs. nucleation undercooling in delta iron of various oxides, nitrides and carbides, (from [5]). 22 Figure 11. Cooling curves measured at different withdrawal rates, (from [34]). 24 Figure 12. Temperature and solubility dependence of the saturation solubility of T i N . [from 62]. 28 Figure 13. Nitrogen distribution in alloy IN718. (from [63]). 29 Figure 14. Size distribution of T i N inclusions as a function of nitrogen content in alloy IN718, with a constant L S T . (from [63]). 30 Figure 15. Dependence of k ™ on Cr content in some superalloys. (from [63]). 30 Figure 16. Temperature dependence of the saturation solubility of T i N in IN718 (from [1]). 31 Figure 17. Relation between superheat and the equilibrium nitrogen contents required for the formation of T i N particles in alloy IN718. 38 Figure 18. Segregation of Titanium in the mushy region in alloy IN718. (from [68]). 39 vii Figure 19. The D S Q apparatus. 44 Figure 20. Schematic of the D S Q apparatus. 45 Figure 21. Temperature distribution along the induction coil in the D S Q (at 2.5 kW). 46 Figure 22 (a) and (b). Exterior of a directionally solidified ingot (directly taken out from the DSQ) . 46 Figure 23. Quenched interface perpendicular to the direction of withdrawal. 49 Figure 24. Quenched interface at an angle to the direction of withdrawal. 49 Figure 25. Schematic quantitative analysis of T i N in (a) D S Q samples, (b) D S samples. 50 Figure 26. A typical T i N inclusion showing a prominent alumina core. 50 Figure 27. Dendrite arm spacings for several directionally solidified superalloys as a function of the cooling rate, (from [69]). 51 Figure 28. Schematic of the mushy interdendritic region in alloy IN718 with T i N Precipitates. 51 Figure 29. S E M picture of the PDS in sample 2(64 ppm [N], D S Q sample). 52 Figure 30. S E M picture of the PDS in sample 1 (68 ppm [N], D S Q sample). 52 Figure 31. Schematic of PDS determination at different sections of the D S cast ingot (sample 5, 64 ppm [N]). 53 Figure 32. Schematic of the stages involved in preparing samples 3 (48 ppm [N]) and 4 (90 ppm [N]). (a) severing the sample in half (b). drilling a hole at the top of the bottom half and, (c). filling the hole with C r 2 N and closing it with the upper half. 59 Figure 33. Schematic of the procedure to prepare homogeneous IN718 samples. 59 Figure 34. T i N distribution in the directionally solidified and quenched samples 1 (68 ppm [N], 2 (64 ppm [N]), and 3 (48 ppm [N]) at a withdrawal rate (R) = 8.1 urn/sec. 61 Figure 35. Temperature distribution of T i N particles in the directionally solidified and quenched samples 1(68 ppm [N], 2 (64 ppm [N]), and 3 (48 ppm [N]) at a withdrawal rate (R) = 8.1 urn/sec. 62 Figure 36. T i N distribution in the Directionally solidified sample 4 (90 ppm [N]) at R = 8.1 um/sec. 63 Figure 37. T i N distribution in the D S samples 5 and 6 ( 64 ppm [N], from the same alloy B) at R = 8.1 urn/sec. 64 viii Figure 38. Nitrogen analysis for the D S cast sample 5 (64 ppm [N] at R = 8.1 pm/sec). 65 Figure 39. T i N distribution in the 90 ppm [N] content D S cast IN718 ingots. Sample 7 at R = 8.1 pm/sec and sample 8 at R = 2.4 pm/sec. 67 Figure 40. Nitrogen distribution in 90 ppm [N] content IN718 ingots. Sample 7 at R = 8.1 pm/sec and sample 8 at R = 2.4 pm/sec. 67 Figure 41. Nitrogen distribution in the 90 ppm [N] content IN718 ingots directionally solidified at 100 microns nitrogen pressure. Sample 9 and 10 at R = 8.1 pm/sec and sample 11 and 12 at R = 2.4 urn/sec. 69 Figure 42. T i N distribution in the 90 ppm [N] content IN718 samples directionally solidified at 100 microns nitrogen pressure. Sample 9 at R = 8.1 pm/sec and sample 11 at R = 2.4 urn/sec. 70 Figure 43. (a) and (b) Etched sections from D S samples (90 ppm [N]) showing T i N particles of the very top of the ingot. 71 Figure 44. (a), (b), and (c) T i N particles on the upper outer surface of the ingot. 72 Figure 45. Nitrogen distribution along sample 13 (90 ppm [N] content) directionally solidified at 400 microns nitrogen pressure at a withdrawal rate of 8.1 pm/sec. 73 Figure 46. T i N distribution along sample 13 (90 ppm [N] content) directionally solidified at 400 microns nitrogen pressure at a withdrawal rate of 8.1 pm/sec. 73 Figure 47. (a), (b), and (c) S E M pictures of the top section of sample 13 showing Large T i N particles. 74 Figure 48. (a) and (b) T i N clusters on the outer surface of the D S cast sample 13 (90 ppm [N]). 75 Figure 49. Nitrogen distribution along the 20 ppm [N] directionally solidified ingots at a withdrawal rate of 2.4 pm/sec. Sample 14 and 15 were D S at a nitrogen pressure of 40 microns and samples 16 and 17 were D S at a nitrogen pressure of 10 microns. 77 Figure 50. T i N distribution along the 20 ppm [N] directionally solidified ingots at a withdrawal rate of 2.4 pm/sec. Sample 14 and 15 were D S at a nitrogen pressure of 40 microns and samples 16 and 17 were D S at a nitrogen pressure of 10 microns. 77 Figure 51. (a), (b), and (c) T i N particles found at the very top of the D S ingot (20 ppm [N], at 40 microns nitrogen pressure). 79 Figure 52. Distribution of T i N along the D S cast IN718 ingots at a withdrawal rate of 2.4 pm/sec. 84 ix Figure 53. Distribution of nitrogen along the D S cast IN718 ingots at a withdrawal rate of 2.4 urn/sec. 85 Figure 54. Distribution of T i N along the D S cast IN718 ingots at a withdrawal rate of 8.1 um/sec. 86 Figure 55. Distribution of nitrogen along the D S cast IN718 ingots at a withdrawal rate of 8.1 umsec. 87 Figure 56. Equilibrium nitrogen contents required for the formation of T i N in the 1/s mushy region at different temperatures, accounting for Ti segregation. 88 Figure 57. (a) Initial stage during D S (dendrites are beginning to form) (b) intermediate Stage, and (c) final stage of D S . 89 Figure 58. Nitrogen distribution along the D S samples at various k values for nitrogen. 91 Figure 59. T i N particle distribution along the directionally solidified IN718 ingots. Data are from the following samples: 90 ppm [N] samples D S cast at 100 microns nitrogen pressure - sample 11 (at R = 2.4um/sec) and 20 ppm [N] content sample 14 directionally solidified at a nitrogen pressure of 40 microns, at R = 2.4 u.m/sec. 93 Figure 60. T i N saturation solubility ahead of the liquidus from the present studies And a typical casting liquidus[63]. 96 Figure 61. (a), and (b) a hexagonal, and a cubic T i N precipitate respectively. 110 Figure 62. (a) and (b) T i N agglomerations bridged by alumina particles. 112 Figure 63. T i N inclusions bridged by alumina particles. 112 Figure 64. Nitrogen distribution in samples 18, 19 ( D S cast under argon pressure of 100 um), 7 and 8 (DS cast at 3.5 x 10"5 Torr, and at 8.1 and 2.4 um/ sec respectively). 116 Figure 65. T i N distribution in samples (18), (19) ( D S cast under argon pressure of 100 u), (7) and (8) (DS cast at 3.5 x 10"5 Torr, and at 8.1 and 2.4 u/ sec respectively). 116 Figure 66. (a) and (b). T i N inclusions on the outer walls of the Directionally Solidified samples 18 and 19. 117 Figure 67. S E M picture of Niobium-carbides. 124 Figure 68. Atypical carbonitride (TiN with an alumina core and surrounded by Nb carbide phase). 125 Figure 69. (a) Delta phase, (b) T i N inclusion precipitated on needle shaped delta phase (c) and (d) show T i N inclusion on Nb rich Laves phases. 126 x Figure 70. Chromium rich deposit on the top of the ingot. 128 Figure 71. A T i N particle amidst chromium globules on the upper outer surface O f the D S cast ingot. 128 xi List Of Tables Table 1. Composition of alloys SX-1 ( From [5]), A M I , M A R - M 2 0 0 [49] and SC-16. [6]. 13 Table 2. Prediction of undercooling of T i N and FffN in Ni-based superalloys. 23 Table 3. Saturation solubility of T i N in IN718 (from [62]). 32 Table 4. Relation between P N 2 and [%N] for alloy IN718 (at T H q = 1340 °C). 36 Table 5. Equilibrium nitrogen contents required for the formation of T i N particles above T l i q in alloy IN718. 37 Table 6. Effect of T i segregation on the equilibrium nitrogen contents required for the formation of T i N below the T H q in alloy IN718. 40 Table 7. Alloy compositions used in the experiments. 42 Table 8. Primary dendrite arm spacing of sample 1 and 2 as calculated from the S E M pictures. 53 Table 9. Average primary dendrite arm spacing at different portions of the D S sample 5 (64 ppm [N]). 54 Table 10. Velocity of rise (V s) of T i N particles of various sizes. 56 Table 11. IN718 Samples used for different experiments. 58 Table 12. Nitrogen distribution along the D S cast ingots a), high temperature gradient b). low temperature gradient. 113 Table 13. Directional solidification experiments conducted under argon atmosphere. 115 Table 14. Effects of H f addition on the equilibrium nitrogen content for H f N formation a tT l i q (1340 o C)inal loyrN718 . 120 Table 15. Relation between the equilibrium nitrogen contents required for the formation of T i N and P N 2 for alloy C M S X - 4 (at 1390°C) . 122 Table 16. Solubility of T i N in alloy C M S X - 4 above T U q ( 1 3 9 0 ° C ) . 122 xn List Of Symbols A Atomic mass (of the element) C E T Columnar to equiaxed transition Co Equilibrium solute concentration (wt%) Ci Composition of the solute element i (wt%) C s Solid solute concentration (wt%) Ci Liquid solute concentration (wt%) D Diffusion coefficient (cm2/sec) D S Directional solidification D S Q Directional solidification and quenching d s Distance travelled by the secondary dendrite arm (cm) dp length of the platform (of the turbine blade) (cm) f Activity coefficient (of the solute element) fs Fraction solidified I Nucleation rate (/sec/cc) G Thermal gradient (°C/cm) A G 0 Standard free energy (Joules) AFT Enthalpy of formation (Joules) k Equilibrium partition coefficient K Equilibrium constant (for the reaction) mi Liquidus slope of a solute element ( °C/wt%) N 0 Nucleation site density (/cc) P c Solutal Pectal number xiii PDS Primary dendrite arm spacing (pm) R Withdrawal rate (pm/sec) AS° Standard Entropy of formation (Joules/°C) S X Single crystal Tii q Liquidus temperature (°C) Tsoi Solidus temperature (°C) A T Undercooling (°C) A T S Dendrite tip undercooling (°C) A T n u c Undercooling for heterogeneous nucleation (°C) T i N Titanium nitride (precipitate) V s Stokes' velocity of rise (for the T i N particle) (pm/sec) *F A parabolic function describing the transverse variation of temperature with distance. e Wagner interaction parameter (for the element) 5 percentage lattice mismatch between the substrate and the nucleant p Density of solid T i N particle (g/cm3) a Density of liquid IN718 alloy (g/cm3) n Coefficient of viscosity (of liquid IN718 alloy) (Ns/m 2) xiv Acknowledgements I would like to thank my supervisor, Dr. Alec Mitchell for his invaluable guidance throughout this work. I would also like to thank Dr. Steve Cockroft, Dr. Warren Poole, and Dr. Desmond Tromans for their useful discussions and suggestions. The help rendered by A. Schmalz in designing, and building of the experimental equipment is duly acknowledged. A special thanks to Philippe Auburtin, for all his help throughout this thesis. I am very thankful to Special Metals Corporation in providing help with the alloys used in this study. I am very grateful to my family for their encouragement and support. Finally, I am very thankful to Mani for making all this possible. xv Chapter 1: Introduction The demands for the operation of turbine engines at increased level of temperatures and reduced engine weight have been the principal contributing factor for the development of nickel base superalloy castings with enhanced mechanical properties. In spite of the improvements, the materials are still found to undergo premature failure at defects such as inclusions. It has been established that the low cycle fatigue crack (LCF) origins are primarily at the clusters of inclusions such as nitrides and only in very clean alloys is failure produced at the carbides or structural features such as grain boundaries [1], as illustrated in Figure 1. From the critical crack size versus inclusion size graph it can be seen that the agglomerate T i N is very detrimental to the mechanical properties. It is further known that the primary carbides and nitrides block the interdendritic fluid flow and thus have an effect on the formation of defects such as microporosity and freckles [2]. Among the non-metallic inclusions, the most important one is titanium nitride (TiN). It is found that the particles themselves not only detoriate the mechanical properties, but also influence the solidification structure [3]. Hence, it is important to know the thermodynamics of formation of nitride inclusions and the solubility of T i N in nickel base superalloys in order to predict under what conditions they may form. Chapter 2 of this report deals with this aspect. A n added problem in superalloys is that, they all have a long freezing range (solidus and liquidus temperatures) owing to their complex compositions. Therefore, they exhibit heavy segregation on solidification which has an influence on the nucleation conditions required for the C E T (columnar to equiaxed transition). Random grain growth has been one of the major problems in the commercial casting of single crystal components for turbine engines. The reasons for the formation of random grains have been attributed to factors like discrepancy in 1 solidification due the complex geometry of the cast component, formation of freckle chains and nucleation etc. It has been shown that the segregation created ahead of the solidifying dendrite tip and the undercooling in this region are contributing factors in the process of nucleation. It is to be noted that under normal industrial conditions of directional solidification in alloys, the undercooling mainly refers to the solutal undercooling and not the thermal undercooling. Solutal undercooling occurs due to the fact that as alloys solidify, they reject solute into the liquid at the interface. The accumulation of solute ahead of the interface effectively results in a region of supersaturation ahead of the dendrite tip. This supersaturation (or the related solute undercooling) represents the driving force for the diffusion of solute at the dendrite tip in an alloy. This undercooling is essential for heterogeneous nucleation to take place. A detailed study on the kinetics of nucleation and random grain formation in superalloys is presented in chapter 2 of this report. However, in practice, secondary grains are found in many single crystal castings without any freckling or any drastic change in the thermal conditions. Hence, the source for the formation of secondary grains in single crystal castings is yet to be established. Though there is no strong evidence to substantiate it, there is a high probability that if the non-metallic inclusions like nitrides are rejected into the liquid immediately ahead of the liquidus isotherm during solidification, and if the heterogeneous nuclei are stable in this region, they are very attractive as nucleation sites. This may be an influential factor in random grain production in single crystal casting. Chapter 4 and 5 of this thesis present the theoretical and experimental studies on the problem stated above. The results and discussions are presented in chapter 6 and chapter 7 respectively, followed by conclusions with the recommendations for future work presented in the final chapter. 2 Figure 1. Fatigue properties of PN718, illustrating the influence of inclusions on premature failure at T = 350 °C; strain amplitude = 0.62% (from [1]). 3 Chapter 2 : Literature review This chapter presents a literature review on the random grain formation and the nucleation aspects of inclusion formation in superalloys. A n account of random grain formation through the columnar to equiaxed transition (CET) in superalloys, and the various explanations and contributing factors which have been attributed to the formation of random grains in superalloys in the past years is presented under the first sections of this chapter. The absence of grain boundaries in the S X materials, and elimination of grain boundary strengthening elements like C , B, Zr allows higher temperature solution cycles and greater alloy microstructural homogeneity which improves alloy creep strength. S X / D S casting alloys are now extensively used for turbine blade materials (both aero and industrial turbine blades). In spite of the desirable mechanical properties, these materials are found to exhibit random grain formation during solidification. It is proposed in this thesis that the T i N inclusions may be a possible source for spurious nucleation in the S X / D S alloys. Since data for S X alloys such as a) the saturation solubility of T i N , and b) thermodynamic data pertaining to nitrogen solubility and T i N formation are not available, alloy IN718 was selected as a model system. Other reasons for selecting alloy 1N718 are: a) it is a very common Ni-based superalloy, and b). studies on alloy IN718 have been done earlier, which are relevant to the present studies. Hence, it makes IN718 alloy a convenient model system. Spurious grain formation in N i - based superalloys, particularly the studies carried out by Pollock and Murphy [4,5] on the formation of equiaxed grains in single crystal superalloys and that carried out by Cockroft et. al [6] in the determination of a criterion for spurious grain formation in superalloy turbine blades, is described in Sections 2.2 and 2.3. A n account of the effects of inclusions in random grain formation in iron alloys and in 4 superalloys, and the theory of heterogeneous nucleation based on the Turnbull-Vonnegut [7] lattice disregistry model are presented in Section 2.4. Following this analysis, an account of the chemistry behind the solubility of nitrogen and the formation of T i N in Ni-based superalloys is discussed under Section 2.5. Sub-section 2.5.3 presents some earlier studies made on the formation of T i N inclusions in alloy IN718. 2.1 Random Grain Formation Owing to their complex compositions, superalloys have a long freezing range. Therefore, they are subjected to heavy segregation on solidification and thus have an influence on the conditions required for the C E T (columnar to equiaxed transition). It has been shown that the segregation ahead of the solidifying primary dendrite tip and the undercooling in this region contribute towards nucleation of equiaxed grains. Though there is no strong evidence to substantiate it, there is a high probability that if the non-metallic inclusions like nitrides are rejected in the liquid immediately ahead of the liquidus isotherm during solidification, they would be very attractive as nucleation sites, which may be influential in random grain formation. 2.1.1 Fundamentals of CET C E T is a sudden change in the grain morphology of a casting from a columnar structure where the grains are oriented parallel to the direction of heat flow, to an equiaxed structure where the grains have a random orientation distribution. The C E T phenomenon involves two processes: 1). formation of equiaxed grains, 2). growth of equiaxed grains and how they hinder the columnar grains. Addressing the first aspect, two main hypotheses have been put forward to explain the formation of equiaxed grains namely, 1). heterogeneous nucleation 2). dendritic fragmentation. The heterogeneous nucleation model has two invariants depending upon the method of casting. 5 In the case of conventional casting the mold is preheated to a temperature above the liquidus temperature of the alloy and the liquid alloy is directly poured into the crucible. The crucible is open at the bottom end which is in direct contact with a water cooled copper chill. According to the "big-bang" model [8], spurious nucleation occurs during the initial stages of solidification due to the drastic cooling of the liquid alloy when it comes into contact with the water cooled copper chill. Alternatively, in the case of directional solidification or single crystal casting, where the mold is withdrawn through a temperature gradient at a specific withdrawal rate, nucleation can occur at anytime during solidification due to supercooling [9]. Studies by Fredriksson et. al [10,11] on lead and aluminum alloys and on steel ingots revealed that the C E T at the vertical solidification front was caused by the development of free crystals. It was proposed that free crystals are formed in the liquid ahead of the solidification front, which is cooler than the liquid in the central part of the mould. When they multiply to large quantities, they physically adhere to the solidification front and thus hinder the growth of the columnar grains. According to this theory, higher the cooling rate of the mould, and lower the solidification process, easier was the C E T . Studies have revealed the occurrence of C E T in experimental ingots depending upon the imposed conditions. For example, the presence of heterogeneous catalysts has caused nucleation during solidification [12]; equiaxed grains have been induced in E S R (electroslag remelted) iron alloys by addition of T i C [13] etc. However, in the case of larger conventional ingots, the most important process to cause the C E T is likely to be the 'big bang' process. The dendritic fragmentation model also has many invariants like the dendrite arm melt-off due to thermal effects such as convection [14]; or fragmentation due to mechanical damage [15]; or due to fall-out from the upper surface of the ingot [16]. Regarding the second aspect of C E T namely the growth of equiaxed grains and how they hinder the growth of the columnar grains, much work has been done. It was established by 6 Doherty et. al., [17] that in order to grow, the latent heat of the growing nuclei should be conducted out to the ingot mold, past the tips of the columnar dendrites. Therefore, the equiaxed grains must grow at a higher temperature than the columnar grains. Witzke et. al., [18] derived the thermal and chemical fields ahead of the growing tips of the columnar dendrites from theories of natural convection and have suggested two necessary conditions for the growth of equiaxed grains: 1). the intensity of constitutional undercooling, and 2). The volume of the undercooled liquid both must reach a sufficient value. This aspect of undercooling occurring ahead of the growing dendrites is discussed in detail in the following sub-section. 2.1.1.1 Dendrite Undercooling Two different types of dendrite undercooling have been identified. One occurs with dendrite growth into an imposed temperature gradient, where the undercooling increases for faster dendritic growth (cellular or constrained dendritic growth) [19,20]. The situation for the columnar or directional solidification is often referred to as constrained growth, where the rate of advance of the isotherms constrains the dendrites to grow at a given velocity. The second where the dendrite growth occurs with no temperature gradient in which, the undercooling increases as the dendrites grow faster [21-23]. The role of temperature in gradient causing this difference was suggested by Doherty et. al., [22] and was verified experimentally and theoretically by Burden and Hunt [19]. The phenomenon of dendrite tip growth has been extensively studied for binary alloys [24-27]. Coates et. al [27] have mathematically extended (with some modifications) this to ternary and multiconstituent alloys and have shown that each solute element makes an individual contribution towards the constitutional supercooling of the interface. Bobadilla [28] and Rappaz et al. [29] have modified the K G T [30] model and have shown that the individual undercoolings 7 of each solute element ' i ' may be summed to determine the overall undercooling of the dendrite tip as: m i = slope of the liquidus Co = equilibrium composition Cj = composition of solute element i ki = distribution coefficient for the element i I(Pc) = O = Ivanstov function of the solutal Pectal number Pc, given by: [31] where, R = rate of advance of the dendrite tip P = radius of the tip Di = liquid diffusion coefficient Equation (1) was further simplified by assuming that the liquid diffusion coefficients for each element are equal , which implies that the Pectal numbers and supersaturations associated with the solutes are equal. Hence, the equation reduces to the expression 3 from which the dendrite tip undercooling for a multiconstituent alloy can be calculated if the growth rate is known. (1) where, (3) 8 for which, the "pseudo-binary" variables C , m , and k are given by: (4) IXc;,, m = ^—= c (5) and S r n , . C * 1 . k i k = ^ — = m.c ( 6 ) Therefore, it can be inferred from equation 1, 2 and 3 that the rate of solidification plays an important role in determining the undercooling of the dendrite tips. 2.2 Spurious Grain Formation In Ni-Based Superalloys Random grains or spurious grains have been observed in single crystal and directionally solidified turbine blade castings, which severely detoriate the mechanical properties. The source for heterogeneous nucleation and undercooling are the primary factors causing C E T and spurious grain formation in Ni-based superalloy castings. Many models have been put forward to explain the phenomenon of breakdown of columnar structure and the formation of equiaxed grains in N i -based superalloys. These theories have considered either of the two views for grain nucleation namely, a). Thermosolutal convection, and b). Constitutional undercooling, which are discussed in the following sub-sections. 2.2.1 Thermosolutal Convection Pollock and Murphy [4,5] have studied grain defect formation in a number of superalloys and their observations have suggested that in high refractory content superalloys, the breakdown of 9 single crystal solidification and the formation of mis-oriented grains and defects such as freckles are dependent on thermosolutal convection processes. It was found that the transition from single crystal to equiaxed grain structure occurred gradually for a fixed alloy composition and withdrawal rate. A s the thermal gradients were decreased, isolated and highly mis-oriented columnar grains appeared until the advance of the single crystal front was completely blocked. A major factor that affected this transition was the primary dendrite arm spacing (PDS). Figure 2 shows the number of grains observed as a function of primary dendrite arm spacing for each of the experiments. a '3 250 ° 200 J •a | 150 -1 o i <» o u l> a a Z loo A 50 0 I Directional Dendritic 1 • 1 ' 1 • 1 n Isolated Misoriented Grains ' • • i i 1 1 1 1 m Equiaxed b o | 1 T T J I T T | I I i j I I 600 800 0 200 400 Primary dendrite arm spacing (u) 11 • ' I 1 1000 Figure 2. The three regimes of grain defect formation and the dependence on dendrite arm spacing, (from [5]). It can be seen from Figure 2 that on the basis of the primary dendrite arm spacing, three regimes of behavior can be identified. 1). Region of single crystal solidification for PDS < 320 um 2). Region of solidification interrupted by occasional nucleation resulting in the formation of mis-oriented columnar grains for a PDS between 320-600 um and 3). Region of equiaxed 10 solidification with complete blockage of the solidification front beyond a PDS value of 600 pm. C E T was generally observed to occur in the third regime. Nucleation and growth of isolated columnar grains and freckles occurred in the intermediate range. A n example of a mis-oriented isolated columnar grain is presented in Figure 3. A systematic study on these defects has revealed that the nucleation on inclusions in the melt cannot be attributed to the formation of these defects, but occasional internal grains with no surface connections were observed. Figure 3. Isolated mis-oriented columnar grains Figure 4. Complete breakdown of single during directional solidification of alloy SX-1 solidification and blockage of solidification (after [5]). Front by equiaxed grains, (after [6]). In experiments where the thermal gradients were decreased below 15 °C/cm, there was a transition from single crystal to equiaxed grains. Figure 4 presents such a transition from the airfoil region to the base of the turbine blade. The complete breakdown of the single crystal solidification has occurred here. In addition to the formation of columnar and equiaxed grains, freckle type defects have also been identified. Figure 5 shows an example of a freckle chain. These are observed in the 11 low-gradient regime. During the unidirectional solidification of the alloys, due to solute rejection there is density inversion developed in the interdendritic liquid, resulting in convective flow and formation of solute-rich channels. Figure 5. A typical freckle chain observed during solidification of alloy SX-1. (from [6]). Pollock [5] has observed that the elements which influenced defect formation to the greatest extent were refractory elements like Re, W, and Ta, which are high density elements (they have densities about two times that of nickel in the solid state) as well as more strongly partitioning elements. It was also observed that freckles were rich in elements that have a partition coefficient < 1. While elements like A l (k=0.98) and Ta (k=0.53) segregated to the interdendritic liquid, there was a strong depletion of elements with partition coefficients > 1 like Re ( k=l .65 ) and W (k=l.37 ) relative to the base material from the freckled regions. Figure 6 shows the relation between the number of freckled chains and the primary dendrite arm spacing. For alloy SX-1, the onset of freckles (as shown in Figure 6) was observed at the same primary dendrite arm spacing 12 c •a sz U •u 3 30 25 20 15 10 Single Crystal I I Columnar Grains and Freckles • QS • • a a Qa rnrnrrnrJIhrrn i n , - i — T — r - i — r -m Equiaxed 100 200 300 400 500 600 700 Dendrite Arm Spacing (microns) 800 Figure 6. The dependence of freckle chain formation on primary dendrite arm spacing, (after [6]). \ A l l o y SX-1 A M I M A R - M200 SC-16 elemerms-^ wt% w t % w t % w t % N i Bal. Bal. Bal. Bal Cr 4.5 7.52 8.4 16 Co 12.5 6.64 9.52 -M o - 2.06 0.03 3 A l 6 5.34 4.98 3.5 T i - 1.25 1.92 3.5 W 5.8 5.59 11.9 -T a 7 7.8 0.10 3.5 N b - 0.01 0.86 -H f 0.16 0.004 1.18 -Re 6.3 - - -C - 80 ppm 0.14 -B - 3 ppm 150 ppm -Zr - 10 ppm 400 ppm -Table 1. Composition of alloys SX-1 (From [5]), A M I , M A R - M 2 0 0 (from [49]) and SC-16 (from [6]). 13 where columnar grains began to nucleate [as shown in Figure 2]. The general composition of alloy SX-1 is given in Table 1. Since the columnar grain nucleation and freckles were found to occur under similar thermal and solutal conditions, it was proposed that convection and dendritic fragmentation contribute to the nucleation of columnar grains and freckle type defects, and ultimately to the formation of random or equiaxed grains. This is not substantiated by the model configuration of remelting times for fragments by Erdeljac et al [32]. According to their model, larger dendrite-like particles (> 1.6 mm) would not melt completely and would be entrapped in the mushy zone. 2.2.2 Constitutional Undercooling The best-suited model to apply to Ni-based superalloys is the Hunt's [33] model for steady-state analysis of columnar and equiaxed growth. According to this model heterogeneous nucleation in the constitutionally undercooled liquid at the columnar front is the only source for equiaxed grains. C E T will occur once the volume fraction of the equiaxed grains ahead of the columnar dendrites reaches a value of 0.49. The critical thermal gradient in the liquid ahead of the columnar dendrites, below which a C E T will occur, is given by: G < 0.617 N 0 1/3 (AT;) 3 A T C (7) Where, G = critical thermal gradient at the liquid/solid interface below which C E T occurs N 0 = nucleation site density A T C = columnar dendrite tip undercooling A T N = heterogeneous nucleation undercooling at which 1 site/cc becomes active 14 Since the solidification rate (R) is incorporated in the equation 7 for undercooling (AT C), both the factors G and R affect the C E T phenomenon. The cooling rate (G x R) during the solidification determines the amount of time that the equiaxed grains have to grow before the arrival of the columnar front. If the equiaxed grains have not grown to a full volume fraction of 0.49, C E T will not occur. Therefore, C E T is favored by high solidification rate, which maximize the solidification front undercooling. The role of the thermal gradient is to provide a finite volume of constitutionally undercooled liquid ahead of the interface in which active nucleation sites may be found. Therefore, C E T is favored by low thermal gradients that maximize the volume of the undercooled liquid. 2.3 Criterion For Spurious Grain Nucleation The formation of spurious grains was investigated by Cockroft et. al [6] in Ni-based superalloy turbine blades of C M S X - 4 using a modified Hunt's model for C E T in A l - C u alloys. A general mathematical model was developed to avoid spurious nucleation in the body of the blade and the extremities of the head and the root. This was done by determining the limiting conditions for the solidification rate 'R ' and the axial temperature gradient ' G ' , which are the most important solidification parameters in a S X or D S casting process. It has been established that the probability of nucleation of equiaxed grains depends upon the dendrite tip undercooling, which can be calculated for binary alloys using the K G T model. Since superalloys are complex multi-component systems, the K G T model was extended, as suggested by Bobadilla [28] and Rappaz [29] and the pseudo-binary values were calculated using equations 3 to 6. The density of equiaxed grains (dn) at a given undercooling d(AT) was calculated using a continuous nucleation distribution of the form of a Gaussian distribution, originally proposed by Rappaz [34]. This is given by: 15 n(AT) = [ . " m a x exp rAT-ATV AT, v J d(AT) (8) where, n m a x , is the total number of nucleation sites available (2.5 x 10s m"3), AIT and ATa (1.5 °C) are the distribution parameters at the center and the standard deviation respectively. Using these values for alloy C M S X - 4 , the value of A T was found to be 6.7 °C. It was found that C E T can occur if the volume fraction of equiaxed grains ahead of the columnar dendrite front exceeded a critical value of 0.66. This value was employed to study the spurious grain formation in the body of the turbine blades. When a columnar dendritic front (which is also the curvature of the liquidus isotherm) passes through extreme enlargements in cross-sections of the component (for e.g. transition from the blade to the outer shroud), depending upon the undercooling occurring during the process two phenomenon can occur [35], as illustrated in Figure 7. shroud. low undercooling capacity liquid 'liq shell mold liquid y — T / x T | i q y/(J undercooled ijsolidls , / V \ m e l t high undercooling capacity formation ot rapid lateral growth of spurious grains secondary dendrite arms Figure 7. Depending upon the undercooling capacity of the given alloy a macroscopic concave curvature of the liquidus isotherm can cause the formation of stray grains (a) or result in fast lateral growth of secondary dendrites (b). (from [35]). 16 Figure 8. Schematic diagram showing the growth of a secondary dendrite arm into a platform in the presence of a transverse temperature gradient. (from[6]). 1) . A concave liquidus isotherm can result in isolated, thermally undercooled regions of the melt. Depending upon the number of nuclei available in the melt and the extent of undercooling in this region heterogeneous nucleation and the formation of spurious grains can occur close to or at the ceramic mold as showed in Figure 7(a). 2) . If nucleation is suppressed, secondary dendrites will grow rapidly laterally towards the far corners of the shroud, as soon as the primary dendrites reach the shroud. This is schematically represented in Figure 7(b). The secondary dendrite arm must grow up to the new lateral surface of the mold in order to propagate the single crystal in these open regions of the liquid by tertiary branching mechanism, as represented in Figure 8 [6]. Once inside this region and unconstrained by neighboring solute fields, the tertiary arms can break free to become new primary dendrites. The rate at which the secondary dendrites grow into the cavity is finite and is a function of the 17 undercooling of the liquid immediately ahead of the dendrite tip. At appropriate conditions, the undercooling adjacent to the mold wall at the extremity of the cavity may increase to such a point when the heterogeneous site would become active prior to the arrival of the secondary dendrite arm. This results in the formation of "stray crystals" [36-38]. Conditions under which spurious grains can form in the turbine blade platform were derived based on assessing whether the secondary arm was able to traverse the length dp of the platform, before the dendrite tip undercooling AT^ exceeded the critical undercooling for heterogeneous nucleation ATnuc [6]. The condition for the formation of spurious grains is given by: A T * ^ A 7 ™ c atds = dp • (9) and the condition under which spurious grains will not form is given by: A T s < ATnuc a t d s = dp (10) where, ds is the distance traveled by the secondary dendrite arm. The undercooling at the tip of the secondary arm is given by: ATs=V.G.t + V(ds) (11) where, V is the mold withdrawal velocity, G is the axial temperature gradient, t is the time, and *F is some function describing the transverse variation in temperature with distance. The distance traveled by the secondary arm is given by: 18 ds(t) = jv,dt 0 (12) where, is the growth velocity of the secondary dendrites. The effect of transverse temperature gradient on the limiting conditions for G and V for spurious grain nucleation in the platform have been studied by varying the temperature gradient G [6]. The parabolic function to describe the relation between the distance traveled by the secondary dendrite arm (ds) and the temperature is given by: [d„ +dm -d~) (13) where, T m a x is the maximum temperature difference between the center of the blade and outside of the mold wall at the extremity of the platform, dp is the platform length, and dm is the thickness of the mould. 10 JL_1'. ~ y 0-1 Equiaxed | Spurious CET mm/s ; I ! <T . . . ; Tmax = 0(Flat) Tmax = 2(C) 0.0O1 / Columnar i Tmax= 5(C) 04001 J 0.00001 _ J ._ . _.. | 0.0001 0401 041 0.1 1 G(C/mm) 10 Figure 9. Effect of transverse temperature gradient in alloy CMSX-4. (from [6]). 19 The results are presented in Figure 9, which indicate that in the platform-limited region, the limiting conditions for V are decreased as the transverse temperature gradient G is increased. This brings out the need for minimizing the extent to which the extremities of the platform are cooler than the body of the blade. Hence, it would be beneficial to keep the extremities of the platform hotter than the body of the blade. This could be achieved by slowing the rate of withdrawal V and the rate at which the radiation shield is opened at the platforms if it could be determined, on-line, exactly when to make the change. 2.4 Inclusions and Random Grain Formation Investigations of superalloys using the E B button melting technique have showed the presence of oxides and nitrides [39] and also the presence of T i N in high temperatures in liquid superalloys [40,41]. Eliminating the source responsible for the production of equiaxed grains in liquid superalloys would also eliminate the production of spurious grains. Studies on revert alloys have attributed the presence of high [N] to be the source for equiaxed grains [42] and also T i N particles were found in revert alloys. Painter et al. [43] found C r N and T i N to be efficient grain refiners in equiaxed castings. Studies on IN738 by L i u et. al. [44] have revealed that additions of B, N i A l , and N b C favored C E T by increasing the size of the equiaxed zone and decreasing the grain size. Pardo [45] found C E T to occur in ferritic stainless steel G T A welds only in the presence of T i N particles, which were also found to be effective nucleating agents at an undercooling of 1.7 °C by Bramfitt [46]. Recently, The effectiveness of T i N as an inoculant for ferritic, and ferrite/austenite stainless steels was studied by Poole et. al [47]. Their studies have revealed that T i N is an effective nucleant for ferritic stainless steels, which have been slowly directionally solidified, but is not an effective inoculant for austenitic stainless steels. This is because, the lattice mismatch between T i N , which has an F C C crystal structure, and ferritic matrix (2.9315 A 0 ) is less when compared to the lattice mismatch between 20 T i N and the austenitic matrix (3.6773 A 0 ) . Hence the presence of T i N induces heterogeneous nucleation in ferritic stainless steel. The concept of lattice disregistry and heterogeneous nucleation is explained in detail below. The Turnbull-Vonnegut model [7] has emphasized the contribution of the crystallographic lattice disregistry between the substrate and the nucleated solid towards heterogeneous nucleation. This model was valid up to a disregistry of 20% (a value above this implies poor coherency between the two phases), and only for substrates and nucleating solid having similar crystal structures. The degree of supercooling of a liquid was postulated to be a parabolic function of the disregistry factor. The expression for undercooling is given by: ATI = 0.556 v 8 y (14) where, A 7; = undercooling 8 = percent lattice disregistry given by: 5 = ^ x 100 % (15) where an is the lattice parameter of the solid inclusion, and Aao is the difference in the lattice parameter between the substrate and the solid inclusion. It is to be noted that the undercoolings thus calculated give an approximate idea as to whether the inclusions can act as effective heterogeneous nucleation sites in the substrate. Bramfit [46] extended this linear disregistry theory to describe similar as well as dissimilar lattice arrangements between the substrate and the nucleating phase. Bramfit [46] and Ohashi [48] have determined the undercoolings before 21 nucleation on a number of nitrides, carbides and oxides in 5-iron, as presented in Figure 10. It is evident from the Figure that TiN is one of the most effective nucleants in delta iron. u o C o J=l -a ' § • 13 .3 § •a o oo .5 "o o T3 e 35 Tumbull-Vonnegut Calculation 30 * Bramfitt Experimental 1 • so, • Ohashi et al. Experimental 25 20 vvc / 15 / Z/C 10 5 0 5 10 IS 20 Lattice Misfit (%) Figure 10. Lattice mis-match vs. nucleation undercooling in delta iron of various oxides, nitrides and carbides (from [5]). Experiments conducted by Fernihough [49] on Ni-based superalloys confirm that heterogeneous nucleation in the region of undercooled liquid ahead of the dendritic tip, due to solute rejection, is the source for equiaxed grain formation. The appearance of spurious grains in the platform of turbine blades can be explained on the basis of whether sufficient volume of undercooled liquid is present and also to the extent of undercooling. The contribution of lattice disregistry towards undercooling can be extended to Ni-based superalloys. Taking into account the lattice expansion at the liquidus temperatures, we can calculate the undercooling required for nucleation on TiN and HfN (which are the stable nitride found in single crystal alloys) in Ni-based superalloys. The values are presented in Table 2. 22 inclusion Lattice parameter at 1340 °C (ao) A 0 % lattice mismatch ( with matrix in N i -based superalloys) Undercooling (°C) ffi2") ATC = 0.556 T i N 4.303 16.7 19 H f N 4.593 21.8 33 Table 2. Prediction of undercooling of T i N and H f N in Ni-based superalloys. Since T i N and N b C have a similar crystal structure (both are cubic) and the lattice parameters are very similar, they can dissolve in one another i.e. they are isomorphous. The observed carbonitrides in IN718 alloys are (Nb, Ti) (C,N). However, in single crystal superalloys for turbine blades containing Hf, the stable carbonitride phase would to be (Hf,Ti) (C,N). The H f / T i ratio may vary between 0.3 - 0.4 in the nitride core, and between 12-19 in the outer carbide phase. A s discussed before, for heterogeneous nucleation to occur, there should be a good lattice fit between the nucleant atom and the matrix atoms. In Ni-based alloys, the precipitated phases invariably have a cubic structure, which has a good lattice fit with the F C C matrix. T i N and H f N have an F C C structure, but the mismatch in the lattice parameter between H f N and the N i matrix is much larger (21.7%) when compared to the mismatch between T i N and the N i matrix (16.8%). Consequently, the undercooling required for nucleation on H f N (calculated using equation 14) is large. This indicates that T i N is a more effective nucleant. In casting single crystal superalloy turbine blades, as the mold is withdrawn into the cold zone, the platforms are observed to supercool by several degrees before the arrival of the dendrites from the main body of the blade. Hence, the contribution of the dendrite arms towards the formation of equiaxed grains can be ruled out. Studies by Fernihough [49] on alloys A M I and M A R - M 2 0 0 have revealed an 23 undercooling of about 3.5 to 3.8 °C at a withdrawal rate of about 5.8 x 10"3 mm/min. O n the other hand, studies by Meyer et. al [35] on turbine blades of SC-16 (composition is given in Table 1) have revealed an undercooling of about 10 K in the leading edge and an undercooling of about 12K in the trailing edge of the blade which also lead to the formation of spurious grains in these regions. Also, the withdrawal rate was found to affect the undercooling produced. B y using a withdrawal rate of 3 mm/min, undercoolings of up to 40K were measured in the outer regions of the shroud. Also, a withdrawal rate of 2 mm/min was found to decrease the undercooling in the shroud area up to 2 K . It is to be noted that in these studies the withdrawal rates used are much higher than the withdrawal rates used in practice for turbine blade casting (< about 1 mm/min), which places a practical limitation on these manipulations. Figure 11 presents the cooling curves measured in the trailing edge at different withdrawal velocities. It can be seen that raising the velocity causes an increase of the cooling rate and of the extent of undercooling. Figure 11. Cooling curves measured at different withdrawal rates [35]. 24 Comparing the calculated values for undercooling of T i N and H f N in Ni-based superalloys with the undercoolings produced during the directional solidification of single crystal turbine blades [38], compels the idea that the source for spurious grain formation may be a foreign substrate like a T i N inclusion. The following sections deal with the thermochemistry behind the dissolution of nitrogen and the formation of T i N in nickel based superalloys. 2.5 Thermodynamics Of Inclusion Formation 2.5.1 Nitrogen in Superalloys Studies on revert superalloys have attributed the presence of microporosity and detoriation in mechanical properties to an increase in nitrogen content [50-53]. R . E Painter et. al. [51] considered that microporosity was not a function of the total nitrogen content, but the form in which the nitrogen was present in the superalloy. The nitrogen content of an alloy comes from the impurity levels of the raw material and contamination during the melting and casting processes. For a specific commercial application, the appropriate nitrogen content is specified in superalloys. Alloy IN718, which finds general application as forged rotating parts, has a specified nitrogen content of 60ppm - 80ppm. The casting alloys, when used for critical application have a lower value of nitrogen content ranging from 15ppm to 25ppm and when used for less stressed parts have a nitrogen content of 40ppm - 80ppm [1]. The reason for this difference is attributed to the formation of casting defects associated in alloys with high nitrogen contents. In the case of single crystal alloys, due to the higher titanium content the solubility of titanium nitride is lower in these alloys. 2.5.1.1 Solubility of Nitrogen The solubility of nitrogen in iron, iron-nickel alloys [54-59, 40] and in multicomponent nickel-based alloys [60] is reported in the literature. On the basis of statistical theory for the lattice 25 model of solutions a formula was derived for the solubility of nitrogen in alloys of iron and nickel, with other transition metals. This is given by: ( n V [%N] = [%N] M e 4 ^ 1 - 7 Z £>J (16) Where, [%N] and [%N] M e are the solubility of nitrogen ( i n mass fractions) at PN2 = 1 atm, and the temperature 1600 °C in an alloy M , and in a pure metallic solvent Me, respectively. X j is the mole fraction of the j-th component in the alloy; gN is the Wagner interaction parameter of nitrogen with the j-th component in the solvent M e ; n is the number of components in the alloy; A M = AMeXMe + ^ A J X J ' where A l is the atomic mass of the j-th component. Using this method, the possibility of calculating the solubility of nitrogen in nickel-based alloys using only the first-order interaction parameters was shown by Bol'shov et al. [60]. Studies have shown that Sievert's law holds for molten, pure iron over modest nitrogen pressures (<10 MPa) and for all titanium bearing alloys up to a nitrogen partial pressure below 1 atm. [40,61]. The dissolution of nitrogen in liquid iron alloy can be written as: ^•N2(g) = [N] (1 wt% in liquid iron) (17) where, k is the equilibrium constant for equation 17 and fN is the activity coefficient of nitrogen in 1 wt% standard state iron. The activity coefficient of nitrogen can be calculated using the known thermodynamics and the following relationship: 26 i o g f N = ^e'N[%i] (19) Where, e'N is the first order interaction coefficient between nitrogen and element i . The standard free energy for the reaction 17 is given by: where, A H 0 and A S 0 represent the standard heat and entropy of nitrogen solution at temperature T. The above equations 17-20 can be applied for nickel-based alloys, assuming a similar behavior, and hence the solubility of nitrogen in Ni-based superalloys can be calculated. 2.5.2 Chemistry of T i N Precipitation It is found that in nickel-based alloys containing niobium, the precipitated nitrides are pure titanium nitrides and do not contain any detectable N b or carbon [1]. Nitrides other than T i N have not been reported in alloys containing Nb. The reaction between T i and N is fully determined for alloys that contain over 0.2 wt% T i . The reaction for the formation of T i N for iron alloys can be applied to Ni-based superalloys and can be written as: A G 0 = A H 0 - TAS° = - R T In k (20) T i N ( s ) = [Ti] (i) + [N] ( 1 ) (21) The Gibbs Standard free energy for the above reaction is given by: A G 0 = - R T In (K s p ) (22) where, K s p is the solubility product of T i N in a superalloy which is defined as: 27 K s p = [Ti wt% x fT i] [ N wt% x fN] (23) Where, fi is the activity coefficient of component i . The activity coefficient of T i is given by the expression : i o g f T i = (24) where, e'TI The first order coefficient of T i and the element i. For practical convenience, Cockroft et. al [62] have incorporated the product of the activity coefficients in the solubility constant and hence re-defined reaction 23 as: K',,, = [ T i w t % ] [ N w t % ] (25) It is evident that equation 21 is temperature controlled such that the equilibrium shifts to the right with increasing temperature as shown in Figure 12. It is also composition dependent, being controlled by the activity coefficients and contents of T i and N , and hence being a function of segregation during freezing. -0.02 0.0 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.16 LOG [ N % I Figure 12. Temperature and solubility dependence of the saturation solubility of T i N . [from 62] 28 2.5.3 Studies in alloy IN718 The solubility of N and T i N in nickel-based alloy IN718 by equilibration methods has been extensively studied [1, 62-64]. The solubility of nitrogen homogenized to the liquidus temperature of alloy IN718 presented given in Figure 13. It can be seen that the solubility of N increases with an increase in superheat. A L L O Y 1NCONEL 718 Liquidui g 1353 deg C ; * 1 1 1 1 1 r - 1 TL 10 20 30 40 SO 60 70 Superheat (deg C) Figure 13. Nitrogen solubility in alloy IN718. (from [63]). A n increase in nitrogen content in the liquid alloy results in the precipitation of large numbers of T i N particles at the liquidus temperature in alloys like IN718, which have low titanium levels. Also, the nitrogen content in the alloy is directly found to affect the precipitated T i N particle size. Figure 14 gives the relationship between the size distribution of agglomerated and individual T i N inclusions and the nitrogen content in alloy LN718, with a constant local solidification time. It can be seen that as the nitrogen content increases, the nitride inclusion size also increases. 29 KO 120 100 80 Total N 60 Content lppn0a) 20 i i i 1 AGGLOKCRAnON —I—' - I -4-- I -'. W . INDMOUAL •4- -f* L INCLUSIONS 0 10 20 30 « ) 50 60 A/erage size of 5 largest nitrides in AO fields [ym) Figure 14. Size distribution of T i N inclusions as a function of nitrogen content in alloy IN718, with a constant L S T . (from [63]). a Dependence of kT 1 N T on Cr cont«i^ / - L t VI * u 1> 1 D Figure 15. Dependence of k ™ on Cr content in some superalloys (from [63]). 30 It has been shown that the saturation solubility of T i N is not only a function of the T i and N content, but can also be affected by Cr content. Figure 15 presents the relation between solubility of T i N and Cr content for various nickel-based alloys. The solubility of T i N is found to increase with Cr content. A s mentioned earlier, the solubility of T i N is a strong function of temperature. This is presented in Figure 16. It can be seen from Figure 16 that the solubility in the interdendritic regions (II) is less than that of the bulk alloy (I). This is due to the segregation of titanium and nitrogen in the interdendritic region (this is discussed in detail under section 4.2). 0 L i 1500 1600 1700 1800 1900 Kelvin, (K) Figure 16. Temperature dependence of the saturation solubility of T i N in IN718 (from [1]). Table 3 presents the saturation solubility relationships for T i N as a function of temperature and partial pressure of nitrogen ( P N 2 )in alloy IN718 [62]. It can be seen from the table that T i N is highly stable at high process temperatures (1700 K) and under high vacuum. The solubility of N in alloy IN718 at T H q (1340 °C) as determined from the solubility product of T i N is found to be about 39 ppm [N]. This implies that 31 the T i N inclusion content in the liquid alloy will become zero, when the equilibrium nitrogen content is less than 39 ppm at Th q . T P N 2 Saturation [N] (K) (atm) (ppm) 1700 4.8 x l O " 5 39 1800 4.0 x l O " 4 112 1900 2.7 x l O " 3 290 2000 1.2 x l O " 2 620 Table 3 Saturation solubility of T i N in IN718. (from [62]). 2.6 Summary Studies on random grain formation in superalloys in the past years have suggested that one or more of the following factors were responsible for the formation of spurious grains in superalloys. They are a). Thermosolutal convection (where the limiting factors are determined to be the temperature gradient and the primary dendrite arm spacing between the advancing dendrites), b). Constitutional undercooling (where the criterion for spurious grain formation are determined to be the temperature gradient, and the solidification rate). On the basis of the Turnbull-Vonnegut lattice disregistry model, T i N has been found to be an effective inoculant in ferritic stainless steels. Application of this disregistry theory to T i N precipitation in Ni-based superalloys reveals that the lattice mis-match between T i N particle and the F C C N i matrix is < 20%. Hence, T i N is an effective heterogeneous nucleant. Considering the following facts, a) the high casting temperatures (typically 1500 °C) used in typical casting experiments; b) the stability of T i N particles under these conditions; and c) comparing the calculated values for undercooling of T i N and H f N in Ni-based superalloys with 32 the undercoolings produced during the directional solidification of single crystal turbine blades [38], compels the idea that the source for spurious grain formation may be a foreign substrate like a T i N inclusion. 33 Chapter 3 : Objectives The principal aim of this study is to verify the hypothesis that the T i N particles are stable in the region ahead of the solidification front (just above the liquidus isotherm) during solidification. Other objectives of the study address some critical issues regarding T i N precipitation during directional solidification. A quantitative study on the T i N precipitation kinetics under different conditions (nitrogen partial pressures, withdrawal rates etc.) would enable to determine if the T i N particles precipitated in the solid/liquid mushy region can float out of the solidifying interface. At this point, it is essential to distinguish between the survival of equilibrium T i N particles that are precipitated just above the liquidus isotherm and the flotation (if at all) exhibited by T i N particles precipitated in the solid/liquid mushy region. Considering the density difference between the T i N particles (5.2 g/cm3) and that of liquid IN718 alloy (7.6 g/cm ), and also taking the effect of buoyancy forces in the liquid alloy, is it reasonable to assume that the T i N inclusions would float out of the solidifying interface? Also, it is to be noted that high temperature gradients cause re-solutioning of T i N particles. Hence, finally the conditions that lead to the flotation of T i N particles have to be determined. 34 Chapter 4; Theoretical Methodology For TiN formation In Alloy IN718 The determination of the equilibrium nitrogen concentration for the formation of T i N at various conditions (temperature, nitrogen pressure etc.) is a crucial and an essential part of this thesis. There are insufficient thermodynamic data or citations available in the earlier studies to theoretically determine the equilibrium nitrogen content required for the formation of T i N in alloy IN718. This chapter describes in detail the theoretical methodology and thermodynamic calculations employed to determine the equilibrium conditions for the solubility of nitrogen and the formation of T i N in alloy IN718. 4.1 Solubility of Nitrogen in alloy IN718 The dissolution of nitrogen in liquid iron alloys can be applied to nickel-based superalloys, on the assumption of a similar behavior, and is represented by: V2N2 <"> [N] (lwt% liquid iron) (26) The free energy of the reaction is given by [65]: A G 0 = 3598 + 23.89T Joules (27) The available values for the first order interaction coefficients of nitrogen with various elements in steel, as determined by Elliot [66], were substituted in equation 19 in order to determine the activity coefficient of nitrogen (fN) in alloy IN718. This value was determined to be 0.065. The solubility constant for the reaction can be calculated using equation 20, and at the liquidus 35 temperature for alloy IN718 (1340 °C) , it was found to be 0.0432. be calculated at various pressures using equation 28. fN[%N] = K ( ^ ) The solubility of nitrogen can (28) This data is presented in Table 4. Nitrogen pressure Units used in experiments P N 2 (atm) [N] ( p p m ) 400 microns 5.5 x 10"4 156 100 microns 1.4 x 10"4 78 40 microns 5.5 x 10"5 49 10 microns 1.4 x 10"5 25 3.5 x IO"5 Torr 4.5 x 10"8 1.4 Table 4. Relation between P N 2 and [%N] for alloy IN718 (at T H q = 1340 °C) From Table 4, it can be seen that for alloy IN718, the solubility of nitrogen is quite high at pressures of about 10"4 to 3.5 x 10"5 atm. 4.2 Solubility of TiN in alloy IN718 The dissolution of T i N in iron alloys can be applied to nickel based alloy systems and is given by equation 29. The free energy for the reaction is given by equation 30[67]: TiN(s) = [Ti](l wt% in liquid iron) + [N](l wt% in liquid iron) (29) A G 0 = 308780 - 114.35T Joules (30) 36 The first order interaction coefficients for T i and nitrogen were determined from the corresponding data available for steel [66], and were assumed to be the same for alloy IN718. This was then substituted in equation 24. The activity coefficients for T i and nitrogen were determined to be 0.372 and 0.065 respectively. The solubility constant K was calculated using equation 22 and was found to be 94.26 x 10"6. The equilibrium nitrogen content required for the formation of T i N can be calculated using equation 31. (%N) = ^ (31) (%Ti)(fTi.fN) Table 5 presents the variation of the equilibrium nitrogen contents with temperature (above the liquidus temperature 1340 °C) Figure 17 is a graphical presentation of the relation between superheat and the equilibrium nitrogen contents required for the formation of T i N , as presented in Table 5. T ° C Superheat ( ° C ) [N] ppm 1340 0 39 1390 50 78 1440 100 149 1490 150 311 Table 5. Equilibrium nitrogen contents required for the formation of T i N particles above Tn q in alloy IN718. Below the liquidus temperature, the segregation effects of T i and nitrogen in the mushy region, along with the effect of temperature play a significant role in reducing the T i N solubility 37 considerably. The segregation of titanium in the mushy region in alloy IN718 is known and is presented in Figure 18 [68]. Though it is established that T i segregation and nitrogen segregation take place simultaneously, the extent of nitrogen segregation is unknown. Table 6 presents the variation of the equilibrium nitrogen contents required for the formation of T i N particles below the liquidus temperature in alloy IN718, taking into account T i segregation. 3 5 0 -i 3 0 0 -2 5 0 -pm) 2 0 0 -z 1 5 0 -1 0 0 -5 0 -0 -0 5 0 1 0 0 1 5 0 2 0 0 superheat (°C) Figure 17. Relation between superheat and the equilibrium nitrogen contents required for the formation of T i N particles in alloy IN718. These results indicate that T i N formation in nickel based superalloys is dependent upon the following factors: a) Temperature, b) nitrogen pressure, c) segregation effects of elements (e.g. T i and nitrogen in alloy IN718), and d) the equilibrium nitrogen concentration. 38 Therefore, the above calculations indicate that when the nitrogen content of the alloy at the liquidus temperature (1340 °C) is below the equilibrium nitrogen concentration of 39 ppm, there is no T i N precipitation. Figure 18. Segregation of Titanium in the mushy region in alloy IN718 (from [68]). 39 T ° C [wt. %Ti] [N] ppm with Ti segregation [N] ppm without Ti segregation 1340 1 39 39 1320 1.5 19 29 1300 1.6 14 22 1280 1.7 9 16 1240 2 4 9 Table 6. Effect of T i segregation on the equilibrium nitrogen contents required for the formation of T i N below the T H q in alloy IN718. 40 Chapter 5: Experimental Studies S X / D S casting alloys are now extensively used for turbine blade materials (both aero and industrial turbine blades). In spite of the desirable mechanical properties, these materials are found to exhibit random grain formation during solidification. It is proposed in this thesis that the T i N inclusions may be a possible source for spurious nucleation in the S X / D S alloys. Since data for S X alloys such as: a) the saturation solubility of T i N , b) thermodynamic data pertaining to nitrogen solubility and T i N formation etc. are not available, alloy IN718 was selected as a model system. Other reasons for selecting alloy IN718 are: a) it is a very common Ni-based superalloy, and b). studies on alloy IN718 have been done earlier, which are relevant to the present studies. Hence, it makes IN718 alloy a convenient model system. 5.1 Alloy Compositions Used For experiments Three different composition A (with 67 ppm [N]), B (with 64 ppm [N]) and C (with 4 ppm [N]) of the nickel-based alloy 718, as given in Table 6 were chosen for the present studies. The equilibrium freezing range for alloy IN718 is 1260 °C (solidus temperature) - 1340 °C (liquidus temperature). 5.2 Sample Preparation The samples to be directionally solidified were machined to small rods of about 10-11 g in weight, 6-7 cm in length and 5.2 mm in diameter, in order to fit inside a high purity alumina crucible plugged with a high purity, low-porosity, non-reactive alumina cement. The samples were thoroughly cleaned and dried with alcohol to remove any machine debris on them before being placed in the crucible. 41 3 u ON 0 0 O C N p OH 0 0 0 S O O U l N m O p (N O p O p <+H X 0 d 1 1 OQ ( N 0 0 C N O O O O < 0.45 r-vn Os rr H 0.99 m °! VO OS 03 H 0 0 t-~ 0 p 0 p 5.27 vri 5.03 > 0.03 C N p rr 0 p 0.07 VO p 0 0 o 2.97 2.94 2.96 u rr 00 »-H 00 0 od O U 0.39 VO C N 0 p (H u 17.9 18.0 17.9 0.07 vo O c-0 p Mn 0.07 p O O 0 0 0 O O O VO 0 p U * S O C N m O © so 0 p O -H -0 /i OS Os I N O in i s / 5 / P .2 s/ s <? • < 1 B < J e S •c cu O , X CD I •a u w C o CN r f 00 o I o o >» o Mg cs (S 0 t -VO 3 T O _ a E / c a. / E 7 ; •S / £ w / s < CQ u 5.3 Apparatus and Procedure The apparatus used is a variation of the classical Bridgeman furnace. It is a directional solidification and quench (DSQ) vacuum induction furnace. A photograph and a schematic of the furnace are shown in Figure 19 and Figure 20 respectively. The main chamber consists of an outer quartz tube and an inner susceptor tube (alumina tube surrounded by molybdenum foil) into which the crucible containing the sample can be freely lowered or pulled up by means of an A . C synchronous motor (30 rph) and a shaft arrangement. The inner susceptor tube is surrounded by a set of three outer alumina tubes for radiation shielding. The tubes rest on a flat water-cooled copper chill, with a provision at the center to effect the quenching process. The sample was heated by means of an induction coil, which surrounds the outer quartz tube. The induction coil is a water-cooled copper coil, connected to a 5 kW, 450 kHz power supply. The top of the susceptor tube marks the beginning of the 'hot zone', which extends to about 140 mm. The temperature gradient, which is a very important parameter in the solidification process, was measured by a W-3%Re and W-25% Re thermocouple. Different runs were performed to measure the temperature gradient, as presented in Figure 21. The Temperature gradient at the lower portion of the coil from where the ingot will be withdrawn was found to be a constant for each run. The average temperature gradient (G) was found to be 8 x 103 °C/m. Another important parameter is the solidification rate (R) for the process. For the present experiments, the rate at which the sample moves thorough the zone is governed by the speed of the A . C synchronous motor. For the 30 rph motor used in the experiments, the withdrawal rate was found to be 8.1 urn/sec. The entire chamber was maintained at a high vacuum of about 3.5 x 10"5 Torr using a diffusion pump and a mechanical pump. This vacuum was found to be very effective in producing samples with a very clean exterior, i.e., without any scale formation or contamination on the surface. The exterior of such a directionally solidified ingot directly taken out from the 43 furnace are presented in Figures 22 (a) and (b). Figure 19. The D S Q apparatus 44 To A.C. Supply-*.-A.C Motor Release knob forqueirching Quartz tube Induction coil Alumina _=J2. shielding tubes o O Water inlet '—» Water outlet Mo, saphyl cushion Adjustable platform To vacuum pump *• To vacuum pump Figure 20. Schematic of the D S Q apparatus 45 Figure 21. Temperature distribution along the induction coil in the D S Q at 2.5 kW. (a) (b) Figure 22 (a) and (b) Exterior of some directionally solidified ingot (directly taken out from the DSQ) . 46 The crucible was lowered into the susceptor tube to the required depth (at which the temperature is well above the liquidus temperature of the alloy) by mechanically moving the shaft. Then, the chamber was evacuated using the mechanical pump and then switching on to the diffusion pump once the pressure reaches below 200 pm (read from the vacuum gauge). When the chamber pressure reaches about 10"4 Torr, the power is switched on. After setting the desired power level of 2.5 kW, the sample is held for half an hour (by this time the sample is completely liquid) and the motor is switched on. The distance moved by the sample inside the hot zone is read by converting the number of turns the shaft makes, as read by the mechanical counter, into the corresponding linear distance. For a complete directional solidification process, after the sample has completely exited the hot zone, the power is turned off and the sample is allowed to solidify (the only heat losses being that of radiation). When a quenching process is desired, it was realized by turning a small lever from outside, which detaches a molybdenum wire connecting the crucible with the shaft, thus releasing the crucible. A s the crucible drops down it enters the quench zone at the bottom of the chamber beneath the copper chill, which is filled with layers of saphyl and molybdenum sheet to provide cushioning effect for the crucible as it falls. The final directionally solidified (DS) or directionally solidified and quenched (DSQ) ingots are metallographically prepared by the following method. 5.4 Metallographic Analysis The samples were found to be very clean indicating a good vacuum through out the experiment. The ingot was mounted on a resin along the length, and ground to a width of about 5mm in grinding wheels of 60, 120, 180, 320 and 600 grits successively. It was then polished first in a 5 pm diamond and then in a 1 pm diamond cloth polishing wheels. If the sample was directionally 47 solidified and quenched, it was divided into equal areas 10mm in height, above and below the interface, along its entire length. The samples were etched with Kallings' II reagent (2g C u C l 2 + 40 ml HC1 + 40 - 80 ml ethanol). After etching the interface, which is the transition from the coarse dendritic columnar region to a fine structure, in each sample can be clearly distinguished. In most of the samples, this interface is perpendicular to the growth direction as shown in Figure 23, and in some cases, it is at an angle to the growth direction as shown in Figure 24, depending upon the heat withdrawal characteristics. If the sample was a completely directionally solidified one, it was again divided into equal areas of 10mm in height from the top (the final portion of the liquid to solidify) to the bottom, along its entire length. The number of titanium nitride inclusions in each region was physically counted while observing through an optical microscope at a magnification of 800x. A schematic of the analysis in the case of a directional solidification and quenching and a complete directional solidification are presented in Figure 25(a) and (b) respectively. The T i N particles are distinguished from other inclusions like carbides and carbo-nitrides and counted. Under an optical microscope, the nitrides can be distinguished from other precipitated phases by their color. The nitrides usually have a pinkish or yellow-orange color, often with a dark nucleus. A typical T i N particle as observed from a S E M (scanning electron microscope) is presented in Figure 26. A n E D X microprobe analysis reveals that most of the T i N inclusions have an aluminum rich core. The carbides appear as grayish or bluish particles and the carbonitrides are found to have a very distinct pinkish core with a grayish outer region surrounding them. Other phases like the niobium rich Laves and needle shaped delta phase (NisNb) (which are unavoidable in nickel alloys containing Nb) were also observed which are described in detail in Appendix 6. The samples are then observed under the scanning electron microscope (SEM) to measure the T i N particle sizes at different regions on the polished specimen. 48 Figure 23. Quenched interface perpendicular to the direction of withdrawal. Mushy zone -10 Top of ingot 0 !0 20 30 40 r 10 - • 2 0 - • 3 0 _^40 _^50 (a) (b) Figure 25. Schematic quantitative analysis of T i N in (a) D S Q samples, (b) DS samples. Figure 26. A typical T i N inclusion showing a prominent alumina core. 50 5.5 Primary Dendrite Arm Spacing In the directional solidification process, the temperature gradient (G) and the solidification rate (R) are the two important solidification parameters. The relation between the cooling rate (G x R) and the dendrite arm spacing has been investigated for many superalloys [69] and this relation is presented in Figure 27. c — i 1 — m — i 1 — m 1 1 — m 1 1—rr DENDRITE SPACINO (lim) PRIMARY ARMS IEGENO 0 IN-100 O IN713C & IN-718 V IN-738LC • IN792 O MAR-M.246 G NI-9AI..4CR 0 NASAIR-100 - i — I i I I 1 I I I ! i i i i i Cooling rate (G x R), °C/sec Figure 27. Dendrite arm spacings for several directionally solidified superalloys as a function of the cooling rate, (from [69]). Primary dendrite Secondary dendrite TiN particle Figure (28). Schematic of the mushy interdendritic region in alloy IN718 with T i N precipitates. 51 The primary aim is to quantitatively verify whether the T i N particles can float out of the solidifying interface. The finer the dendrites, the greater is the chance that the T i N particles may be caught up between the dendrites. Therefore, for the present studies a coarse dendritic structure i.e. large primary dendrite arm spacing is desired in the ingots. The wide interdendritic spacing offers less resistance to fluid flow, hence could facilitate flotation. Figure 28 is a schematic of the mushy interdendritic region of alloy IN718 along with the precipitated T i N particles. Please note that the typical value for the primary dendrite arm spacing is about 400 microns and that for the T i N particles is about 6 microns. (The figure is magnified approximately 50 times). Figure 29. S E M picture of the PDS in Figure 30. S E M picture of the PDS in sample 2 (64 ppm [N], D S Q sample). sample 1 (68 ppm [N], D S Q sample). The primary dendrite arm spacing was directly measured by observing the etched cross-sections of the sample in the S E M (Scanning electron microscope). Figure 29 and Figure 30 present such an S E M picture of the primary dendrite cross-sections observed in samples 2 and 1 respectively. 52 Table 8 gives the average primary dendrite arm spacing for Samples 1 and 2 as measured from the S E M pictures. Corresponding to the value of (G x R) used in the present study (0.065 °C/s) , the primary dendrite arm spacing (PDS) is found to be about 400 pm from Figure 27, which agrees well with the values obtained here (Table 8). Alloy Sample 1 (68 ppm [N]) Sample 2 (64 ppm [N]) P.D.S (pm) 414 425 Table 8. Primary dendrite arm spacing of sample 1 and 2 as calculated from the S E M pictures. Section PDS (pm) 1 (424) 2 (409) > 3 (413) 4 (410) * 5 (443) Figure 31. Schematic of PDS determination at different sections of the D S cast ingot (sample 5, 64 ppm [N]). Also the DS cast sample 5 (64 ppm [N]) was sectioned horizontally at different portions along its length, and each portion was ground and polished at the flat surface. These polished surfaces were etched with Kallings' reagent to observe the transverse microstructure under the scanning 53 D S 718 ingot electron microscope. Figure 31 presents a schematic of this procedure and Table 9 presents the average PDS at each section, as measured from the S E M pictures. The PDS values at these sections were almost a constant. Section PDS (pm) 1 424 2 409 3 413 4 410 5 443 Table 9. Average primary dendrite arm spacing at different portions of the DS sample 5 (64 ppm [N])-5.6 Buoyancy Effect on T i N Alloy IN718 doesn't exhibit density changes due to solute segregation in the S/L mushy region, unlike the S X / D S alloys which freckle. Note that FN718 has a very low carbon content (0.008 wt%), which may contribute to the fact that the alloy does not exhibit freckling. In the case of S X / D S alloys, freckling occurs by density inversion due to solute rejection in the segregation zone. There is fluid flow in this region and thus thermosolutal convection is established [70-72]. The maximum velocity at the center of a freckle plume was calculated in the NH4CI-H2O castings and in using a modified Poiseuille equation [70,72]. This was later extended to lead alloy systems (Pb-Sn and Pb-Sb) by Sarazin et. al [70] and the freckle velocity was determined to be 100 mm/sec. 54 From a point of view of buoyancy, an inclusion like T i N is similar to a solid particle, which can either sink or float in the liquid depending upon whether its density is higher or lower relative to the density of the liquid alloy, and the viscous drag on it. Since, the density of T i N particles (5.2 g/cm ) is lesser than the density of IN718 (7.6 g/cm ) at the Tn q , in principle, buoyancy forces ought to be the driving force in causing the flotation of the T i N particles. In general, the Stokes' principle can be applied to calculate the "velocity of rise" or "velocity of sink" of a spherical particle in a viscous liquid. For the purpose of the thesis, the Stokes' principle is modified to accommodate cubic T i N particles by replacing the volume of the spherical particles by the volume of the cubic particles (see equation 32-33 for details). Any discrepancy arising from the "shape effects" is assumed to be small and is ignored for the purpose of calculation. Using this modified Stokes' principle, the velocity of rise of cubic T i N particles in the liquid IN718 alloy may be calculated as follows: The viscous drag on the particle of density p, and volume d (assuming that the T i N particles are cubic with each side = d units) in a liquid alloy of density a, and coefficient of viscosity r\, is given by: 67rr|dVs = (mass of [TiN]s ) x g - ( m a s s of equal volume of [IN718] i i q) x g (32) 67rridVs = ( d 3 p ) g - ( d V ) g (33) Vs = ^ ( / Q - c Q g 6nr/ 2 where, g is the acceleration due to gravity ( 9.8 m/s ) (34) 55 Equation 34 is the Stokes' relation. This indicates that the T i N particle floats up with a velocity V s . Also, it can be inferred that due to the buoyancy effect, the larger the particle the higher is the rise velocity. A s an example, substituting the corresponding values of n,p,a, and r\ the velocity of flotation of a typical T i N particle of size d = 6 pm in equation 34 gives: (6xl(r 6) 2(5200-7632) 6x3.14x0.0045 V s x9.8 (35) V s = - 10 urn/sec (36) Using a similar calculation, the velocity of rise of T i N particles of various sizes are tabulated and presented in Table 10. T i N particle size (pm) Stokes' velocity of rise (pm/sec) 4 4.5 5 7 6 10 8 18 10 28 Table 10. Velocity of rise (V s) of T i N particles of various sizes. 56 From Table 10, using the present withdrawal rate of 8.1 um/sec, it is reasonable to expect that in principle T i N particles of size > 6 um should be able float out of the solidifying front of the ingot. In the case of lead alloys systems (Pb-Sn and Pb-Sb) that exhibit freckling, the velocity at the center of the freckle plume is about 100 mm/sec [70]. Assuming a similar value for freckle rise velocity in the case of D S / S X Ni-based alloy systems, there is a high possibility that at such low withdrawal rates, the smaller T i N particles might be carried by the freckle plumes from the mushy zone and thus might be able to "escape" out of the solidifying interface. 5.7 Experiments Basically, the experiments conducted consist of two steps. In the first step, the samples were directionally solidified and quenched (DSQ) in an induction furnace, under vacuum. The quenched samples were then metallographically prepared and a quantitative assessment of the distribution of nitrides in various regions above and below the liquidus was made. This helps in understanding the chemistry of T i N at the interface (in the solid/liquid mushy region, and above the liquidus). The second step consists of a complete directional solidification of the samples (the same IN718 ingots with various nitrogen contents) in the furnace, under different conditions of pressures and withdrawal rates, and assessment of the final D S ingots to see whether the nitrides are collected at the top of the ingot. This step helps to determine what are the conditions that can or cannot lead to the T i N precipitation and their flotation above the solidification front, thus verifying or disproving the assumed hypothesis. Table 11 presents the different samples were used for the D S Q (directional solidification and quenching) and DS (directional solidification), and the conditions under which these experiments were conducted. 57 Sample Alloy [N] content (ppm) Experiment Conducted Pressure Withdrawal rate (um/sec) 1 A 67 DSQ 3.5 x 10"" Torr 8.1 2 B 64 DSQ tc 3 C + Cr2N 48 DSQ tt 4 C+ Cr2N 90 DS tt 5,6 B 64 DS tt 7 " (homogenized) 90 DS " tt 8 90 DS 2.4 9, 10 a 90 DS 100 pm nitrogen pressure 8.1 11, 12 a 90 DS a 2.4 13 a 90 DS 400 pm nitrogen pressure 8.1 14, 15 ii 20 DS 40 pm nitrogen pressure 2.4 16, 17 a 20 DS 10 pm nitrogen pressure 2.4 Table 11. FN718 Samples used for different experiments Sample 3 (48 ppm [N]) and sample 4 (90 ppm [N]) were prepared from the same IN718 alloy C by adding appropriate amounts of C r 2 N . A schematic of this procedure is presented in Figure 32. The ingot was cut in half and C r 2 N was placed on top of the bottom half of the sample. The upper half was then placed over the lower half and the whole ingot was carefully slipped into the alumina crucible. Samples 5, and 6 were prepared by melting the alloy C with appropriate amounts of chromium nitride in a separate induction furnace (in order to obtain homogeneous nitrogen composition in the ingot) and were cast into long cylindrical bars of about 5 mm in diameter. 58 u (a) (b) C r 2 N -> IN718 ingot (c) Figure 32. Schematic of the stages involved in preparing samples 3 (48 ppm [N]) and 4 (90 ppm [N]). (a) severing the sample in half (b). drilling a hole at the top of the bottom half and, (c). filling the hole with C r 2 N and closing it with the upper half. O O O N i plate O M o foil -4-O Induction coil ^ _ 0 O p. Crucible >• C r 2 N IN718 pieces O Q • Alumina cement Figure 33. Schematic of the procedure to prepare homogeneous IN718 samples. 59 A schematic of this procedure is presented in Figure 33. These bars were cut to lengths of about 5 - 6 cms and used as individual samples to be directionally solidified in the D S Q furnace. A l l the samples were metallographically prepared and examined for T i N precipitates, as discussed in section 5.4. 5.7.1 Nitrogen Analysis O f The D S Cast Samples A higher concentration of T i N particles in the ingot would produce a higher nitrogen concentration in that area. Hence, if there is flotation of T i N particles and if the T i N particles collected on top of the ingot, the upper portion of the ingot should register a higher nitrogen content. Hence, the different portions of the DS cast samples were cut and the individual pieces were sent out to Special Metals Corporation for nitrogen analysis by a conventional gas analysis technique. 60 Chapter 6; Results 6.1 Directional Solidification And Quenching Experiments A set of directional solidification and quenching (DSQ) experiments were performed on samples 1(67 ppm [N]), 2(64 ppm [N]), 3(48 ppm [N]). After the ingots were taken out of the furnace, they were prepared and polished (as discussed in section 4.6) for quantitative analysis of T i N particles. The number of T i N particles per sq. mm in each region above and below the quenched interface were calculated. From the experimental data, graphs were plotted between the density of T i N (number of nitrides per m m 2 area) and the distance along the length of the sample (with respect to the liquidus in the case of D S Q samples. This is presented in Figure 34. The number of nitrides is averaged over each of the regions i.e. from 0- (-10), 0-10, 10-20, 20-30 mm etc. and 6 — • — samplel —•—sample 2 E 5 —*— sample3 E IT Ifl 4 -in m ;o 'E 3 o ci c 2 1 T -20 -10 0 10 20 30 40 50 (top) Distance from interface (mm) (bottom) Figure 34. T i N distribution in the directionally solidified and quenched samples 1(68 ppm [N], 2 (64 ppm [N]), and 3 (48 ppm [N]) at a withdrawal rate (R) = 8.1 urn/sec. 61 plotted at distances o f . . . -15,-5, 5, 15, 25 mm etc. from the interface (refer to Figure 25). The y-error bars in the graphs are about ± 10-15% . 6 —*— sample 1 5 "••»- sample 2 E E —*— sample 3 ri-te 4 trides - 3 -j of ni 2 Ar d c 1 / / T 1500 1400 1300 1200 1100 1000 (top) Temperature from interface (C) (bottom) Figure 35. Temperature distribution of T i N particles in the directionally solidified and quenched samples 1(68 ppm [N], 2 (64 ppm [N]), and 3 (48 ppm [N]) at a withdrawal rate (R) = 8.1 pm/sec. Figure 35 shows the relation between the temperature distribution and the density of T i N found in the various regions of the ingot. The temperature vs. nitrides graph 35 is plotted by converting the distances in graph 34 to the corresponding temperatures (in °C) from the temperature profile (Figure 21). • In the quenched samples, Figure 34-35, it was found that the number of T i N particles above the T i i q is zero and the nitrides are found only below the Tn q up to a temperature of about 1340 °C. 62 • From Figure 34 it can be seen that sample 3 (48 ppm [N]) shows a large number of T i N particles compared to sample 1 (67 ppm [N]) and sample 2 (64 ppm [N]). 6.2 Directional Solidification (DS) Experiments 6.2.1 Samples 90 ppm [NI, and 64 ppm [NI A high nitrogen content sample with about 90 ppm nitrogen was prepared in a manner described in section 5.7 (Figure 32), and a complete directional solidification was carried out on it. Also, a complete D S was carried out on samples 5 and 6 (64 ppm [N], each from alloy B). Distance from top of ingot (mm) Figure 36. T i N distribution in the directionally solidified sample 4 (90 ppm[N]) at R=8.1 pm/sec. The ingots were metallographically prepared and the density of T i N particles in each region was calculated. From the experimental data graphs were plotted between the density of T i N particles 63 and the distance from the top of the ingot (taken as the reference). These are presented in Figures 36 and 37 for the 90 ppm (sample 4) and the 64 ppm (samples 5,6) respectively. 0 H 1 1 , , r -5 15 25 35 45 55 Distance from ingot top (mm) Figure 37. T i N distribution in the DS samples 5 and 6 ( 64 ppm [N], from the same alloy B) at R = 8.1 um/sec. • In the D S Q Sample 3 (48 ppm [N]) and the D S sample 4 (90 ppm [N]), Figure 34 and 36 respectively, most of the nitrides are found to be concentrated around the central portion of the ingot. • A complete D S on samples 5 and 6, Figure 37 shows that very few nitrides are formed at the top section and virtually no T i N particles are found in the top most portion of the ingot, where the final liquid has solidified. Most of the nitrides are found at a much lower level in the ingot. 64 Figure 38. Nitrogen analysis for the DS cast sample 5 (64 ppm [N] at R = 8.1 urn/sec). • The nitrogen analysis of the DS cast sample 5 (obtained from Special Metals Corp.) is presented in Figure 38. The Y-error bars are about ± 1-2 ppm. It can be seen from the nitrogen analysis graph that there is high nitrogen content at the bottom of the sample rather than the upper portions. Following these experiments, homogeneous IN718 samples were prepared in order to have a uniform [N] distribution through out the sample. These homogenized samples were subsequently used for further directional solidification experiments. 65 6.2.2 Homogeneous Sample Preparation From the earlier results obtained, homogeneous IN718 samples of various nitrogen concentration were prepared in a separate low frequency induction furnace. A schematic of this set up is shown in Figure 33 in section 5.7. The sample is cast into long bars of about 5mm in diameter. A set of high nitrogen content samples (90 ppm [N]) and low nitrogen content samples (20 ppm [N]) were separately cast in the induction furnace. Further DS experiments were carried on these samples at various conditions of pressure and withdrawal rates, as described in the following sub-sections. A l l the directionally solidified samples were metallographically prepared and the T i N particles were quantitatively assessed. 6.2.3 90 ppmfNl sample DS at 3.5 x 10 5 Torr The 90 ppm [N] content homogeneous samples 7 and 8 were then directionally solidified in the D S Q apparatus at a vacuum of 3.5 x 10"5 Torr at two different withdrawal rates of 8.1 um/sec and 2.4 um/sec respectively. a. At a withdrawal rate of 2.4 um/sec A metallographic examination of the polished ingot surface indicates a higher T i N concentration at the lower portion of the ingot and a lower T i N concentration at the ingot top. This is presented in Figure 39. There is a corresponding increase in the nitrogen concentration at the bottom of the ingot as presented in Figure 40. b. At a withdrawal rate of 8.1 um/sec: The T i N distribution graphs and the nitrogen distribution graphs are presented in Figure 38 and Figure 40 respectively. 66 3 . 5 -, Distance from ingot top (mm) Figure 39. T i N distribution in the 90 ppm [N] content DS cast LN718 ingots. Sample 7 at R = 8.1 urn/sec and sample 8 at R = 2.4 urn/sec. 5 15 25 35 45 Distance from ingot top (mm) Figure 40. Nitrogen distribution in 90 ppm [N] content 1N718 ingots. Sample 7 at R = 8.1 pm/sec and sample 8 at R = 2.4 pm/sec. 67 From the figures it can be seen that at the lower withdrawal rate (R=2.4 urn/sec), nitrogen content as well as the T i N distribution is very low at the top of the ingot when compared to the bottom. At the higher withdrawal rate (8.1 um/sec), the nitrogen distribution is almost a constant and there is a slight increase in the T i N density at about 10 to 25 mm from the bottom of the ingot. 6.2.4 90 ppm [NI DS at 100 microns Nitrogen Pressure To study the T i N precipitation and flotation behavior in a high nitrogen content FN718 alloy, homogeneous 90 ppm [N] content samples were selected. Two sets of experiments were conducted at 100 microns nitrogen pressure at two different withdrawal rates (2.4 um/sec and 8.1 um/sec). At a high nitrogen pressure of about 100 microns, the equilibrium nitrogen content was calculated to be 78 ppm (Table 4), which is much higher than the equilibrium nitrogen content for T i N formation (which is 39 ppm). C. At a withdrawal rate of 2.4 um/sec : Samples 11 and 12 were directionally solidified at a withdrawal rate (R) of 2.4 um/sec. The nitrogen analysis data is presented in Figure 41. It can be seen that there is a high nitrogen concentration in the upper portions of the ingot. The T i N distribution along the ingot is presented in Figure 42. Though the metallographic examination of the ingot reveals a higher T i N concentration at the center of the ingot, large T i N particles > 10 microns in size were found clustered at the very top of the ingot, where the final liquid has solidified. This is presented in Figure 43 (a) and (b). Also, the upper outer surface of the ingot reveals large T i N particles (> 10 microns in size) and clusters of T i N particles sticking to it, as presented in Figure 44 (a), (b) and (c). 68 Distance from ingot top (mm) Figure 41. Nitrogen distribution in the 90 ppm [N] content IN718 ingots directionally solidified at 100 microns nitrogen pressure. Sample 9 and 10 at R = 8.1 um/sec and sample 11 and 12 at R = 2.4 um/sec. D . A t a withdrawal rate of 8.1 um/sec : The T i N distribution along the ingot is presented in Figure 42. A similar observation in the T i N distribution was made here as in the case (c). Large T i N particles were found at the very top of the ingot as presented in Figure 43 (a) and (b). The nitrogen distribution along the ingot is presented in Figure 41. From the graph it can be seen that the nitrogen concentration at the upper portions of the ingots is higher than that in the lower portions. 69 1 4-5 15 25 35 Distance from ingot top (mm) 45 Figure 42. TiN distribution in the 90 ppm [N] content IN718 samples directionally solidified at 100 microns nitrogen pressure. Sample 9 at R = 8.1 um/sec and sample 11 at R = 2.4 um/sec. 70 039915 28KV ki'.?bk'\?'.6w (b) Figure 43 (a) and (b) Etched sections from D S samples (90 ppm [N]) showing TiN particles of the very top of the ingot 8 3 9 9 1 4 S 8 K V k'l28K*B.'&UB (a) 71 (b) (c) Figure 44. (a), (b), and (c) T i N particles on the upper outer surface of the ingot. 6.2.5 90 ppm [NI Sample D S at 400 microns Nitrogen Pressure Directional solidification experiments were conducted with sample 13 ( 90 ppm [N] content) at a withdrawal rate of 8.1 um/sec. under a nitrogen pressure of 400 microns. After metallographic analysis of T i N distribution in the sample, it was sent out for nitrogen analysis. At a nitrogen pressure of 400 microns, the equilibrium nitrogen concentration was found to be 156 ppm (Table 4). The nitrogen analysis for the ingot is presented in Figure 45. This figure indicates a high nitrogen concentration at the upper portion and a low nitrogen concentration at the bottom of the ingot. The T i N distribution (presented in Figure 46) is consistent with the nitrogen analysis result. Also, large number of T i N particles of size > 10 microns were found at the very top of the ingot where the final liquid had solidified. 72 2 7 1 3 4 , , , , , 1 , , 5 10 15 20 25 30 35 40 45 Distance from ingot top (mm) Figure 45. Nitrogen distribution along sample 13 (90 ppm [N] content) directionally solidified at 400 microns nitrogen pressure at a withdrawal rate of 8.1 um/sec. 9 i Distance from ingot top (mm) Figure 46. T i N distribution along sample 13 (90 ppm [N] content) directionally solidified at 400 microns nitrogen pressure at a withdrawal rate of 8.1 um/sec. 73 (a) I t 8 3 2 9 1 5 2 8 K V X l . 5 8 K * 2 < 032915 28KV XI130K"£3!Iui (b) (c) Figure 47. (a), (b), and (c) S E M pictures of the top section of sample 13 showing large T i N particles. 74 (b) Figure 48. (a) and (b) T i N clusters on the outer surface of the D S cast sample 13 (90 ppm [N]). 75 This is shown in Figure 47 (a), (b) and (c). A n examination of the outer upper surface of the ingot indicates large T i N particles and T i N clusters sticking to it. Figure 48 (a) and (b) are some S E M pictures of the T i N clusters adhering to the outer wall of the D S cast ingot. 6.2.6 DS Experiments With 20 ppm [NI Samples The 20 ppm sample was selected to understand the T i N precipitation behavior in a low nitrogen content sample. Two arbitrary values for nitrogen pressure were selected 40 microns and 10 microns. The two conditions were selected such that the equilibrium nitrogen contents at these pressures are above and below the equilibrium nitrogen content for T i N formation i.e. 39 ppm. The equilibrium nitrogen concentration at a nitrogen pressure of 40 microns was calculated to be 49 ppm and the equilibrium nitrogen concentration at the nitrogen pressure of 10 microns was calculated to be 24 ppm. (please refer to Table 4). E. At a withdrawal rate of 2.4 um/sec, and at 40 microns nitrogen pressure: Two separate experiments were conducted on FN718 ingots with 20 ppm [N] content. Samples 14 and 15 were directionally solidified at a nitrogen pressure of 40 microns and at a withdrawal rate of 2.4 pm/sec. After metallographic observations, the samples were sent out for nitrogen analysis. Figures 49 and 50 present the nitrogen concentration and the T i N distribution in the ingots. The nitrogen analysis, as well as the T i N distribution indicate a higher concentration at the upper portions of the ingot. Large T i N particles (> about 10 microns) were found at the very top of the ingot as presented in Figure 51 (a), (b), (c) and also on the outer walls of the ingot. 76 Figure 49. Nitrogen distribution along the 20 ppm [N] directionally solidified ingots at a withdrawal rate of 2.4 pm/sec. Sample 14 and 15 were DS at a nitrogen pressure of 40 microns and samples 16 and 17 were DS at a nitrogen pressure of 10 microns. sample 14 sample15 sample 16 sample 17 15 25 35 45 Distance from ingot top (mm) 5 5 Figure 50. T i N distribution along the 20 ppm [N] directionally solidified ingots at a withdrawal rate of 2.4 pm/sec. Sample 14 and 15 were DS at a nitrogen pressure of 40 microns and samples 16 and 17 were DS at a nitrogen pressure of 10 microns. 77 F . A t a withdrawal rate of 2.4 um/sec at a nitrogen pressure of 10 microns: The overall nitrogen concentration and T i N distribution is more or less a constant through out the DS ingot as presented in Figures 49 and 50. N o T i N inclusions were found at the top most portions of the sample and on the outer walls. 832915 20KV XI ! 30K* 2 § ! lutd (b) 78 (C) 832915 28KV X1 .78K 1 7 . S u a Figure 51. (a), (b), and (c) T i N particles found at the very top of the DS ingot (20 ppm [N], at 40 microns nitrogen pressure). 79 Chapter 7: Discussions 7.1 Directional Solidification And Quenching Experiments The results from the directional solidification and quenching (DSQ) experiments performed on samples 1(67 ppm [N]), 2(64 ppm [N]), 3(48 ppm [N]) indicate there is a high probability that a combination of such low pressures and high temperature causes severe degassing of the nitrogen from the ingot. Hence, the upper portions of the ingot are depleted of the nitrogen content. Consequently, there is no T i N formation in these regions. There is also a possibility that the T i N particles dissolve (undergo re-solutioning) at such high temperatures above the liquidus and precipitate at lower temperatures during solidification. The carbon content in the solidifying ingot is found to affect T i N precipitation. While a lower carbon content (.008 wt % in sample 3 (48 ppm [N])) promotes the precipitation of T i N particles, a higher carbon content (.03 wt% in sample 1 (67 ppm [N]) and sample 2 (64 ppm [N])) favors carbonitride formation. In the case of single crystal alloys, since they have very low carbon content one would anticipate the same "pure" nitride formation effect. 7.2 Directional Solidification (DS) Experiments Directional solidification experiments carried out on sample 3(48 ppm [N]) and sample 4 (90 ppm [N]), indicate a higher T i N concentration at the center of the ingot. This is probably due to an experimental artifact that there is a high probability that the C r 2 N had not homogenized (due to a lack of a stirring action in the induction furnace) and had remained around the central region of the ingot where it was placed. From the DS experiments on samples 5 and 6 (64 ppm [N]) content, it is clear that flotation is not affected. If the buoyancy effects did affect T i N flotation as assumed, then in 80 principle T i N particles (at least those > 6pm) should have been found at the top of the ingot, where the final liquid has solidified. Since the samples have a high nitrogen content (about 64 ppm), it needs to be established experimentally that either the T i N are not all at the top, or stuck on the outer surfaces of the sample. Also, i f the T i N particles are indeed influenced by buoyancy forces and obey Stokes' law, a lower withdrawal rate (< 6 um/sec, as determined for a typical T i N particle with a size of about 6 microns) should be able to affect such a flotation. DS Experiments at 3.5 x 10"5 Torr Directional solidification experiments on homogeneous 90 ppm [N] content samples 7 and 8, conducted in a high vacuum, indicate that during solidification a combination of high temperature and low pressure leads to heavy degassing of nitrogen from the upper portions of the ingot. Hence, there is a low nitrogen concentration and a corresponding lower T i N concentration in the upper portions of the ingot. The importance of nitrogen gas in promoting the formation of T i N inclusions is evident from the above results. The theoretical calculation for the equilibrium nitrogen content required for the formation of T i N particles in alloy IN718 (as discussed in chapter 4) has claimed a value of 39 ppm [N] content at the liquidus temperature. At a pressure of 3.5 x 10"5 Torr(4.5 x 10"8atm) and at the liquidus temperature of 1340 °C, the [ N ] e q m falls to a very low value of 1.4 ppm (Table 4). DS Experiments under Nitrogen Pressure In the D S experiments on the 90 ppm [N] content samples 9-12, conducted under a nitrogen pressure of 100 microns nitrogen pressure, the calculated equilibrium nitrogen concentration is about 78 ppm at Tn q . Hence, there is a heavy T i N precipitation. There is flotation of T i N particles that were precipitated in the liquid alloy, as can be seen from Figure 41. Larger T i N particles > 10 microns in size are able to float out and get accumulated on the top of the ingot. This is 81 because, they have a velocity of rise > about 28 um/sec (calculated using the Stokes principle), which is an order higher than the withdrawal rate (2.4 um/sec) used in the experiments. Though the nitrogen content at the very bottom of the samples 9 and 11 is high, the corresponding T i N distribution (from Figure 42) is low. If there were smaller T i N particles (<about 3 microns in size) sticking to the outer walls of the ingot at the bottom, it was not possible to observe them under the S E M . There may be a possibility that very small T i N particles could be sticking to the outer wall of the ingot at the bottom, which could have contributed to the increase in the [N] content. At a nitrogen pressure of 400 microns, the equilibrium nitrogen concentration was found to be 156 ppm (Table 4). Figures 45 and 46, indicate a flotation of T i N particles precipitated in the liquid. Also, at such high nitrogen pressure there is a possibility of nitrogen pick-up from the atmosphere, which could have partially contributed to an increase in the nitrogen content at the very top of the ingot. The 20 ppm sample was selected to understand the T i N precipitation behavior in a low nitrogen content sample. At a nitrogen pressure of 40 microns, the equilibrium nitrogen content is determined to be 49 ppm (at Ti i q ) . In the segregation zone, due to the combined effects of temperature, segregation of T i and nitrogen, T i N solubility decreases. The observed flotation is due to the precipitated equilibrium titanium nitrides in the liquid. A s seen from the Figures 49, and 50 none of the T i N particles precipitated in the solid/liquid mushy region exhibit flotation. At a nitrogen pressure of 10 microns, the calculated equilibrium nitrogen content is 24 ppm. This value is below the equilibrium nitrogen content required for the formation of T i N particles, which is about 39 ppm. Hence, the overall nitrogen concentration and T i N distribution is very low and more or less constant through out the D S ingot as presented in Figures 49 and 50. Due to a decrease in the nitrogen content, there is a corresponding reduction in the T i N precipitation. 82 7.3 Summary of TiN distribution and Nitrogen analysis The nitrogen analysis and T i N distribution for samples 7 through 17 are presented in the following graphs. Figures 52, and 53 present the T i N distribution and nitrogen distribution along the D S cast 718 samples 7 through 17 at withdrawal rate of 2.4 urn/sec respectively. The T i N distribution and nitrogen distribution along these samples at a withdrawal rate of 8.1 um/sec are presented in Figures 54 and 55. Figures 52 and 53 indicate that in samples 11, and 12 (which are 90 ppm [N] content samples, D S at 100 microns nitrogen pressure) there is some flotation of T i N particles precipitated in the liquid. The equilibrium nitrogen content at this pressure is determined to be about 78 ppm at Tn q . Where as, in Figure 53, samples 14 and 15 (which are 20 ppm [N] content samples DS at 40 microns nitrogen pressure), indicate no such flotation exhibited by T i N particles precipitated in the solid/liquid mushy region. A similar T i N flotation effect can be observed in sample 13 (90 ppm [N], DS at 400 microns nitrogen pressure), as presented in Figure 55. The equilibrium nitrogen content at this pressure is determined to be about 156 ppm at Tn q . Though it is evident that there is flotation of T i N particles precipitated in the liquid, there is a high probability that there may be some nitrogen pick-up from the atmosphere at such high nitrogen pressures. Nitrogen Distribution; A n interesting observation from the nitrogen analysis graphs 53 and 55 is that, although the equilibrium nitrogen contents of the samples 11-15 are high, the final nitrogen content is much reduced. The final ingots invariably exhibit a net reduction in nitrogen content. A possible explanation is as follows: There are two significant processes taking place during solidification 83 namely, a), segregation of T i and nitrogen into the bulk liquid. nitrogen (in the solid/liquid mushy region) and b) rejection of T°C Tu, 1340 1320 1300 1280 T s o l 1260 bottom Figure 56. Equilibrium nitrogen contents required for the formation of T i N in the 1/s mushy region at different temperatures, accounting for T i segregation. Figure 56 presents a schematic of IN718 ingot during directional solidification with the corresponding equilibrium nitrogen contents required for the formation of T i N , in the solid/liquid region of the DS ingot. Below the liquidus, the T i N solubility decreases as presented in Table 6 and Figure 56 (due to temperature effects as well as segregation of T i and nitrogen). Above the liquidus temperature, the solubility of T i N increases as presented in Table 5. Assuming that the liquid alloy is stationary (no liquid flow), the diffusion of nitrogen can be given by the equation: Az (37) Where, D = Diffusion coefficient ( « 10" cm /sec). Substituting the value of D in the above equation, the diffusion rate was found to be about 31.6 pm/sec. This value is considerably higher than the present withdrawal rates (2.4 um/sec and 8.1 pm/sec). Therefore, nitrogen is rejected 88 into the bulk liquid during solidification. As solidification proceeds, the bulk liquid becomes richer in nitrogen and there is heavy T i N precipitation just above the liquidus. A schematic of this process is presented in Figure 57 (a). Top bottom E M High superheat^, region T i N precipitates T, i q Nitrogen enriched region Nitrogen rejection, T i and nitrogen segregation Region enriched in Nitrogen and T i N Solubility < 39 ppm (a) (b) (c) Figure 57 (a) Initial stage during D.S (dendrites are beginning to form) (b) intermediate stage, and (c) final stage of D.S. Figure 57 (b) presents a schematic of the solidification process in the intermediate stage, where the dendrites are fully grown. Due to a combination of a) T i and nitrogen segregation and b) rejection of nitrogen, there is a decrease in T i N saturation solubility from the tip to the base of the dendrites in the solid/liquid mushy region. This facilitates T i N precipitation in the mushy region. Above the liquidus, there is T i N precipitation. These particles have sufficient time to grow into larger particles and float. Well above the liquidus, there is dissolution of the T i N particles due to the high superheat in these regions. The T i N particles are re-precipitated later during solidification (at lower temperatures). 89 A s solidification progresses, the bulk liquid is continuously enriched in nitrogen. There is continuous rejection of nitrogen into the atmosphere. Hence, there is a nitrogen "sink" established above the solidifying ingot. The final portion to solidify is enriched in nitrogen and T i N precipitates, as presented in Figure 57 (c). These are the T i N particles that were precipitated above the liquidus during solidification, which had sufficient time to grow into large particles. Below the final portion to solidify, the nitrogen concentration is well below the solubility limit at Tnq (i.e. < 39 ppm). In the process of nitrogen analysis, each portion was ground on all sides and the top in order to determine the nitrogen concentration. There is a high possibility that if the final portion of the ingot is lopped off, the nitrogen-enriched portion may be lost and the analysis shows the nitrogen distribution below the final portion to solidify, which has a lower nitrogen distribution. In this study, the thermodynamic equations describing the solubility of nitrogen, and formation of T i N in iron alloys have been applied to IN718 alloy system. The available data for the first order interaction parameters for nitrogen and titanium with various elements in steel were directly applied to alloy IN718. This data was then used to calculate critical values like the activity coefficient of nitrogen and that of titanium in alloy IN718. Hence, there is a possibility that the initial value for the equilibrium nitrogen content for the formation of T i N thus calculated might be incorrect, which could account for the apparent reduction in nitrogen content observed in the final D S IN718 ingots. The nitrogen concentration at various portions of the ingot during solidification can be calculate using the Scheil equation: C s = kC0(l-fs)k"1 (38) At this point, the phase diagram for nickel-nitrogen is not available in literature and hence the value for the partition coefficient is unknown. To roughly understand the nitrogen distribution along the ingot, the value of partition coefficient for nitrogen was calculated using the physical 90 data from the present experiments. Using the data from sample 11 (90 ppm [N] content sample DS at 100 microns nitrogen pressure), k = C s / C i (39) where, C s = the solute concentration in the solid (10 ppm [N]), Q = solute concentration in the liquid (80 ppm [N]). Hence, the partition coefficient was found to be about 0.125. This value for k was substituted in equation 38. Figure 58 presents the distribution of nitrogen along the ingot. This graph is superimposed on the nitrogen distribution graphs for samples 11 and 12. A s the value of k is decreased, the Scheil equation is found to fit the nitrogen curves determined by the present experiments. Hence, it can be speculated that the partition coefficient of nitrogen in alloy FN718 would be less than 0.1. 0 1 0 2 0 3 0 Distance from ingot top (mm) Figure 58. Nitrogen distribution along the DS samples at various k values for nitrogen. It is to be noted that studies by Bower et. al. [73] on Al-4 .5%Cu alloys have revealed that at sufficiently high thermal gradients and low withdrawal rates (which are similar to the casting 91 conditions used in the present studies) there is significant bulk diffusion of the solute. In this case, the composition vs. fraction solidified is given by the modified Scheil expression : a ak k-\ + \ 1 -k-1 Jc-l (40) V Where a= GiD, /C 0 Rmi Di = diffusion coefficient in liquid Gi = thermal gradient R = withdrawal rate mi = slope of the liquidus (from the phase diagram) Hence, by substituting the casting parameters from the present experiments and the appropriate values for parameters 'a' and 'k'(which at this point are not available), a better fit for nitrogen distribution along the DS ingot could be obtained. To summarize, in-spite of a high equilibrium nitrogen content in samples 11 - 15 the net reduction in the nitrogen content of the final DS ingots may be due to the following reasons: a) segregation of T i and nitrogen in the solid/liquid mushy region, b). rejection of nitrogen into the bulk liquid, c). grinding off the final portion of the ingot, which is enriched in nitrogen and T i N during nitrogen detection, and d). a possible error (if any) in determining the equilibrium nitrogen content required for the formation of T i N in alloy IN718, since the thermodynamic data for nitrogen solubility and T i N formation in iron alloys have been directly applied to IN718 alloy system. 92 7.4 T i N Particle Distribution From Experiments (c) sample 11 (90 ppm sample, at 2.4 pm/sec and 100 microns nitrogen pressure), and (f) sample 14 (20 ppm sample at 2.4 um/sec and 40 microns nitrogen pressure), the T i N particle size distribution were studied along the D S ingots. At various distances from the top of the ingot, the largest T i N particles were selected and their size were measured in the S E M . Figure 59 is a graphical presentation of the T i N particle distribution along the D S ingots. The following observations are made from Figure 59: 30 • sample 11 g 2 5 "f A sample 14 E 20 f v 1 t« 15 i is • I 0 0 10 20 30 40 50 Distance from ingot top (mm) Figure 59. T i N particle distribution along the directionally solidified IN718 ingots. Data are from the following samples: 90 ppm [N] samples DS cast at 100 microns nitrogen pressure -sample 11 (at R = 2.4 um/sec) and 20 ppm [N] content sample 14 directionally solidified at a nitrogen pressure of 40 microns, at R = 2.4 pm/sec. • T i N particles larger than about 9 microns are found at the very top of the ingot. Particles as big as 25 microns were observed in this region. The velocity of rise of the particles is > 28 pm/sec, which is much higher than the withdrawal rates of 8.1 pm/sec or 2.4 pm/sec. Hence 93 it can be concluded that the buoyancy effect does cause flotation of T i N particles greater than at least 10 microns. These large T i N particles are the ones precipitated in the liquid that had sufficient time to grow and which exhibit flotation. From about 5 to 35 mm from the top, the T i N particle size varies from 8 microns to 4 microns such that, the bigger particles are found towards the upper portions and the smaller particles are more towards the lower end. The velocity of rise of these particles ranges from 18 um/sec to 4.5 um/sec respectively (From Table 10). Though there is indication of flotation, it is not much pronounced in this case. At this point it appears that this withdrawal rate is still insufficient to cause a complete flotation of particles < about 8 microns. Below about 35 mm from the top of the ingot, T i N particles < 4 microns are predominant. The velocity of rise of these particles is < 4um/sec, which is very close to the withdrawal rates used in the experiments. Also, the viscous drag on these particles is large and this may be responsible in causing this "sinking" effect. These T i N particles may be the ones that are precipitated later during solidification in the solid/liquid mushy region. 94 Chapter 8; Conclusions and Recommendations 8.1 Summary The principal aim of this work was to study the behavior of T i N particles during the directional solidification in alloy IN718 and establish whether the T i N particles could have any influence on random grain formation. The literature review indicate that random grain formation in many single crystal turbine blades is found to occur even without any drastic changes in the casting conditions, which were earlier determined to be the principal contributing factors towards random grain formation. N o quantitative work has been done in this problem and the relative importance of the source for secondary grains is yet to be established. In casting single crystal superalloy turbine blades, it is found that the platform supercools by several degrees. Experiments conducted on SC-16 turbine blades have indicated that the mold withdrawal rate directly affects the undercooling produced. Undercoolings of up to 40K have been produced in these experiments [38]. According to the Turnbull-Vonnegut theory, a lattice disregistry of < 20% between the matrix and the precipitate is then required for the precipitate to act as a stable heterogeneous nuclei. Taking into account the lattice disregistry between the F C C nickel matrix (in Ni-based superalloys) and the different inclusions, it has been established that the T i N inclusions have a lattice disregistry of about 16.8% and would act as a nuclei at an undercooling of about 19 °C. Hence, T i N particles can be potential heterogeneous nuclei in N i -based superalloys. The D S experiments have concluded that under certain conditions the T i N particles are stable in the liquid alloy ahead of the solidification front (above the liquidus). Therefore, there is a high possibility that the T i N particles can act as potential heterogeneous nuclei for random grain formation in these alloys. 95 1. Conditions Required for the Formation of TiN The equilibrium nitrogen concentration required for the formation of T i N particles in alloy rN718 was calculated to be 39 ppm ( at 1340 °C) using the available thermodynamic data. Figure 60 is a graphical presentation of T i N saturation solubility ahead of the liquidus in the present studies and a typical casting liquidus [63]. 90 30 20 10 20 30 40 Superheat(C) present studies literature 50 60 Figure 60. Variation of the equilibrium nitrogen content required for the formation of T i N with superheat, from the present studies and a typical casting liquidus [63]. Results from the present studies indicate that T i N precipitation will not take place once the equilibrium nitrogen content is below this value of 39 ppm (at Tnq). To provide nuclei, the nitrogen content must exceed the saturation solubility of T i N at Tij q , since we have established that T i N particles precipitated in the solid/liquid mushy zone will not float out of the interface. The nitrogen partial pressure is an important variable in T i N precipitation. Higher PN2 can increase the equilibrium nitrogen content of the alloy to a value adequate in causing T i N 96 precipitation. The required nitrogen partial pressure levels for precipitation of T i N particles in typical castings of alloys is comparable to that of ambient vacuum pressures commonly used. Hence, pressure in the D S furnace is an important industrial variable. T i N precipitation is greatly reduced when the D S experiments are conducted in a vacuum of < 3 x 10"5 Torr. This condition causes nitrogen degassing from the upper portions of the ingot, which in turn reduces the equilibrium nitrogen content for the formation of T i N in these regions. Consequently, T i N precipitation is restricted only to the lower portions of the ingot. This effect may also play a role in the industrial observations of random grain formation. T i N precipitation can be minimized by conducting the D S experiments at a pressure less than 3 x 10"5 Torr or by using a very low nitrogen content sample (< 4ppm). This would ensure that the equilibrium nitrogen content for T i N formation is below the critical value and hence T i N precipitation could be eliminated above Tn q . 2. Segregation and nitrogen rejection: During solidification, two important phenomenon takes place namely, a), segregation of T i and nitrogen in the liquid/solid mushy region, and b). rejection of nitrogen into the bulk liquid. In the solid/liquid mushy zone (the region in between the T s o i i d u s and Tnquidus of the alloy), there is segregation of titanium and nitrogen. The segregation of T i is known, as presented in Table 6. Where as, the segregation of nitrogen is yet to be determined. The segregation effects, combined with the temperature gradient in this zone causes a reduction in the saturation solubility value. Due to a decrease in the saturation solubility, there is T i N precipitation in this region. Along with the segregation effects, there is also nitrogen rejection into the bulk liquid, which increases the nitrogen concentration in the liquid above Tu q . 97 The nitrogen diffusion rate was calculated to be about 31 urn/sec, which is greater than the present withdrawal rates of 2.4 and 8.1 um/sec. Due to nitrogen enrichment in the bulk liquid, there is T i N precipitation in the region above Tu q . The T i N particles thus precipitated can undergo flotation into the regions well above Tn q . Due to a high superheat in these regions, they get dissolved and may be re-precipitated later as solidification progresses. As solidification proceeds, the bulk liquid becomes more and more enriched in nitrogen and T i N precipitates, finally becoming accumulated at the top, which is the final portion of the ingot to solidify. The T i N particles found at the top of the final ingot, are the ones that were precipitated above the Tii q during solidification and which had sufficient time to grow into large particles (>10 microns in size). The rest of the ingot below this nitrogen rich portion has a nitrogen content less than the saturation solubility value of 39 ppm. A combination of the following factors such as: a). T i and nitrogen segregation in the liquid/solid mushy region and b). nitrogen rejection into the bulk liquid, c). the grinding up of the top portion of the ingot (about 2mm from the top) and the sides, and d). a possible error (if any) in determining the equilibrium nitrogen content required for the formation of T i N in alloy IN718 (since the thermodynamic data for nitrogen solubility and T i N formation in iron alloys have been directly applied to IN718 alloy system, assuming a similar behavior), might have contributed towards the apparent reduction in nitrogen contents (to values < 25 ppm ) found in all the DS samples 11-15, irrespective of their initial high equilibrium nitrogen contents. At this point, due to the unavailability of thermodynamic data pertaining to IN718-nitrogen phase diagram, partition coefficient of nitrogen etc, the Scheil equation could not be applied to calculate the nitrogen distribution along the DS ingot. In order to get a rough estimation of the nitrogen distribution, data from the present experiments were used for k (the partition coefficient for nitrogen) and also, different partition coefficient (k) values for nitrogen were employed in order to fit the graphs obtained from Scheil equation to the experimental 98 graphs. A value of 0.07 for k was found to be a reasonable fit. Hence, it is speculated that the partition coefficient for nitrogen may be a value < 0.1. 3. Flotation of T i N particles The experiments conducted with alloy IN718 indicate that the buoyancy effects in the liquid alloy do influence the flotation of T i N particles. Theoretical studies (Section 5.6) on T i N particles have revealed that since the density of T i N particles is lower than the density of the liquid alloy, the T i N particles can float in the liquid alloy depending on the viscous drag on them. For the purpose of this thesis the Stokes' principle, originally applied to spherical particle-viscous liquid system, has been extended to cubic T i N particle-liquid IN718 alloy system to calculate the velocity of rise of the T i N particles. Any discrepancy arising from the "shape effects" is assumed to be small and hence neglected. Present experiments with 718 alloys indicate that using the withdrawal rate of 2.4um/sec, T i N particles of size > 10 microns (which have a rise velocity > 28 um/sec) are able to float and eventually get collected at the top of the ingot. These are the particles that were precipitated above the solidification front (Thq) during DS. Also, at the low withdrawal rates employed (8.1 and 2.4 um/sec), these T i N precipitates had sufficient time to grow into larger particles and hence undergo flotation. O n the other hand, particles of size between 8 and 4 microns are found throughout the ingot with a distribution such that the bigger particles are towards the upper portions and the smaller particles are towards the lower portions. This effect again substantiates the fact that the buoyancy forces do influence T i N flotation, but the effect is more pronounced when the relative difference between the velocity of rise of T i N particles and the withdrawal rate is substantial. These are the T i N particles precipitated in the solid/liquid mushy region. 99 The smaller T i N particles < 4 microns are found at the bottom portions of the ingot which indicates that these particles may be precipitated later during solidification in the solid/liquid mushy zone and thus didn't "escape". Hence, for substantial flotation to occur the withdrawal rate should be at least an order lesser than the velocity of rise of the T i N particles. 4. Low Nitrogen Sample: When the sample has a very low nitrogen content, by directionally solidifying the sample in a nitrogen atmosphere the nitrogen content of the liquid alloy can be increased such that it exceeds the equilibrium nitrogen content required for T i N precipitation. Hence, under such condition, it is possible to have T i N precipitation as well as flotation of bigger T i N particles (>10 microns in size). 8.2 Recommendations For Future Work • The ingots used for the present experiments are about 5 mm in diameter and 50 mm in length, which is very small, compared to the ingots used in practice. Hence, a similar investigation could be carried out by redesigning the D S Q furnace to accommodate a larger ingot. Also, the induction coil design can be manipulated in order to give a shallow temperature gradient with a longer plateau region, to achieve a lesser superheat in the alloy and yet maintain a fairly constant temperature above the liquidus temperature of the alloy. • Since in practice random grain formation has been observed in single crystal turbine blades like in alloy C M S X - 4 and alloy SC-16, it would be beneficial to study the inclusion behavior in these alloys rather than in alloy IN718 (a preliminary theoretical study is presented in Appendix 5. Hence, the present work can be extended to these single crystal alloys. 100 A preliminary theoretical study on H f N formation in superalloys and the effect of H f addition in alloy IN718 has been done in the present studies (presented in Appendix 4). Since H f N is found to be the chief nitride inclusion in single crystal superalloys containing Hf, similar experimental studies can be carried out in alloys containing hafnium. 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A , "Segregation Arising From Natural Convection During Solidification", E P D Congress '91, The M i n . Metals and Materials Soc. Pp. 511-522, 1991. 72. Sarazin. J.R, and Hellawell. A , "Channel Flow in Partly Solidified Alloy Systems", Advances in Phase Transitions, pp. 101-115, 1987. 73. Bower. T .F , Brody. H . D , and Flemings. M . C , "Measurements of Solute Redistribution in Dendritic Solidification", Transactions of The Met. Soc. A I M E , V o l . 236, pp. 624-634, 1966. 108 Appendix: 1 A n interesting aspect regarding the shape of T i N particles in alloy IN718 is that two different shapes are encountered in this study: a), cubic and b). hexagonal. In the previous studies related to T i N precipitation in iron and nickel alloys, only cubic T i N particles have been reported. It is further seen that the hexagonal T i N particles are very often associated with an alumina core. Figure 61 (a), and (b) present a hexagonal and a cubic T i N precipitate respectively, found in the directionally solidified IN718 samples from the present studies. The smaller T i N particles (of size < 4 microns) often are cubic in structure and are found to be devoid of an alumina nucleus. Although, few cubic T i N particles have been observed with an alumina core. (a) 109 (b) Figure 61. (a), and (b) a hexagonal and a cubic T i N precipitate respectively. Hexagonal T i N particles have not yet been reported in literature. The hexagonal T i N particles found in this study are invariably associated with an alumina core. The alumina particle has a hexagonal crystal structure. Since the T i N particle precipitates on the alumina nucleus, this may be influential in causing the T i N particle to grow in a hexagonal crystallographic orientation. It needs to be proved by using conventional techniques for crystal structure determination, like the x-ray crystallographic method i f they are in fact hexagonal T i N particles. 110 Appendix: 2 The T i N agglomerations are very often found to be bridged by alumina particles. This indicates that the alumina particles act as a nucleus for the precipitation of T i N particles and the agglomerate T i N particles are held together by alumina particles. Figure 62(a), and (b) show portions from the top of the 90 ppm [N], 718 DS cast 90 ppm [N] content samples (9) and (13) at (a nitrogen pressure of 100 microns and 400 microns respectively) . The Figures show large number of T i N agglomerations with alumina particles in between holding them together. A n E D X microprobe analysis of these portions shows about 44% alumina and 55% T i (TiN distribution). Figure 63 presents T i N particles which are bridged by alumina particles. These observations substantiate the previous studies by Mitchell et al [63] which have indicated that the T i N agglomerations are bridged by alumina particles. This might be the reason why T i N particles sticks preferentially to the ceramic mold surface in the present study. (a) l l l p. • 1 0 3 9 9 1 4 28KV X£.20K *13.6ui» | (b) Figure 62 (a) and (b) T i N agglomerations bridged by alumina particles. 0 3 9 9 1 5 2 8 K V 55K" 4 ! 3 u i Figure 63. T i N inclusions bridged by alumina particles 112 Appendix : 3 Two directionally solidified IN718 ingots (were prepared in the induction furnace at a vacuum of 10 microns at two different temperature gradients. Sample 8H was directionally solidified at a high temperature gradient of 27 °C/cm and sample 8L was directionally solidified at a low temperature gradient of 11 °C/cm. The withdrawal rate employed was 33 um/sec. These samples were then severed into four equal pieces. A portion from each piece was sent to Special Metals Corp. New York, for nitrogen analysis. This data is presented in Table 12. Sample number (high G) Nitrogen (ppm) 8H1A (top) 42.1 8H2A 34.8 8H3A 39.9 8H4A (bottom) 62.4 Sample number (Low G) Nitrogen (ppm) 8L1A 57.8 8L2A 43.6 8L3A 55.2 8L4A 61.4 (a) (b) Table 12. Nitrogen distribution along the D S cast ingots a) high temperature gradient, b) low temperature gradient. From Table 4, the equilibrium nitrogen content at a pressure of 10 microns is determined to be 25 ppm. In Chapter 7 it is discussed that rejection of nitrogen into the bulk liquid is a very significant factor which contributes to the accumulation of nitrogen and T i N particles at the top portion of the ingot. Also, using equation 37 the nitrogen diffusion rate was calculated to be about 32 um/sec, which is comparable to the withdrawal rate of 33 um/sec at which these 718 samples are cast. Hence, nitrogen diffusion is not fast enough to cause a significant nitrogen 113 degassing which have been observed in samples that were directionally solidified at a lower withdrawal rate of 8.1 or 2.4 um/sec. We still see a decrease in the nitrogen content, but not as much as that observed in the previous cases (in samples 11-15). If there is the same partitioning between the liquid alloy and the solute, we should see a decease in the total nitrogen content, which we do, and there should be an elevated nitrogen content at the top, which there is. The top portion of the ingot was polished and observed under the S E M for T i N particles. T i N particles were not found at the top of the ingot. This is because, at this withdrawal rate, the T i N particles neither have sufficient rise velocity to float, nor sufficient time to grow into larger particles and float. Since the nitrogen analysis shows a nitrogen content above the saturation solubility value (> 39 ppm at the liquidus temperature), there were T i N particles precipitated above the Tu q (1340 °C) , but these particles must have been a part of the solution re-precipitated nitrides. 114 Appendix : 4 Effect of argon: Homogeneous 90 ppm [N] samples were directionally solidified under argon pressure to study the influence of an inert atmosphere on T i N precipitation. Table 13 presents the directional solidification experiments conducted on the 90 ppm [N] content samples 18 and 19. Sample 90 ppm [N] Withdrawal rate R (um/sec) 18 8.1 19 2.4 Table 13. Directional solidification experiments conducted under argon atmosphere. Figure 64 presents the nitrogen distribution and Figure 62 presents the T i N distribution along the D S cast samples 18 and 19 along with that for samples 7 and 8 [90 ppm [N] samples DS cast under vacuum of 3.5 x 10"5 Torr, at 8.1 and 2.4 um/sec respectively] for comparison. • From Fig. (64) it can be seen that the nitrogen contents are almost a constant for experiments conducted under argon pressure, when compared to the experiments conducted under vacuum (which indicate a higher nitrogen content at the bottom of the ingot relative to the upper portions). • From Figure 65 it can be seen that there is a higher T i N concentration at the top of sample 7 (at R= 8.1 pm/sec), when compared to samples 7 and 8 cast under vacuum. A n examination of the outer surface of the samples 18 and 19 indicate large T i N particles sticking to the outer walls (which could not be taken into the nitride counting). This presented in Figure 66 (a) and (b). Hence, the T i N counting is found to be very subjective. sample 18 5 10 15 20 25 30 35 40 45 Distance from ingot top (mm) Figure 64. Nitrogen distribution in samples 18, 19 (DS cast under argon pressure of 100 microns), 7 and 8 (DS cast at 3.5 x 10"5 Torr, and at 8.1 and 2.4 um/sec respectively). 5 10 15 20 25 30 35 40 45 Distance from ingot top (mm) Figure 65. T i N distribution in samples 18, 19 ( DS cast under argon pressure of 100 microns), 7 and 8 (DS cast at 3.5 x 10"5 Torr, and at 8.1 and 2.4 um/sec respectively). 116 Figure 66 (a) and (b). T i N inclusions on the outer walls of the Directionally Solidified sampL 18 and 19. 117 From the above D S experiments conducted on IN718 samples under argon pressure, it was found that there is almost a uniform nitrogen distribution along the ingots, irrespective of the withdrawal rate. Also it appears that during directional solidification, the larger T i N particles stick to the outer surface of the ingot. The purity of the argon gas used for this study is about 99.99%, containing about 15 ppm nitrogen. Hence the presence of nitrogen may have influenced the results to a certain extent. 118 Appendix : 5 Effect of hafnium addition : In single crystal Ni-based superalloys (like alloy C S M X - 4 , where the nominal composition of T i and H f are 1 wt % and 0.1 wt. % respectively), the stable nitrides found are hafnium nitrides. Taking the general reaction for the formation of T i N (equation 21, T i N <=> [Ti] + [N]) an attempt was made here to determine what would be the effect on the solubility of T i N if H f is added to IN718 alloy. Assuming that H f N formation is similar to T i N formation in Ni-based superalloys, the reaction can be written as: H f N s o [Hf], + [N], (41) The solubility constant is given by: K H f N = [%Hf] [%N]. f H f . f N (42) The activity coefficient of H f was assumed to be the same as that for T i (0.372, since no further data were available for Hf). Considering the free energy of formation of T i N (-73 Kcal / mol) and H f N (-81 Kcal / mol) at 1500 K , and substituting them in the free energy equations for reactions 21 and 37, the solubility constant for H f N can be calculated. From this data, the equilibrium nitrogen concentration required for the formation of H f N can be calculated. Considering an alloy with 0% T i , and gradually increasing the H f content, the equilibrium nitrogen content for the formation of H f N has been calculated at the liquidus temperature (1340 °C) . This data is presented in Table 14. It can be seen that with H f addition, the solubility limit decreases. 119 Wt. %Hf [N] e q m ppm 0.4 6.3 0.6 4.2 0.8 3.2 1 2.5 1.2 2.1 1.4 1.8 Table 14. Effect of H f addition on the equilibrium nitrogen contents for H f N formation at Tn q (1340 °C) in alloy IN718. 120 Appendix : 6 Extending the calculations for the solubility of nitrogen and formation of T i N in alloy IN718 (discussed in chapter 2), a similar procedure is adopted here for a typical single crystal alloy C M S X - 4 . Solubility O f Nitrogen: The dissolution of nitrogen in liquid iron alloys can be applied to Ni-based superalloys: ^•N2(g) = [N] (1 wt% liquid iron) (42) The solubility constant (k) for the above reaction is given by: k - m (43) r N i Where, fn is the activity coefficient of nitrogen. The free energy of the reaction is given by [65]: A G 0 = 3598 + 23.89T Joules (44) Applying a similar procedure to calculate the activity coefficient of nitrogen in alloy IN718( as discussed in chapter 2), the activity coefficient of nitrogen in alloy C M S X - 4 was found to be 0.323. The solubility constant for the equation 43 at the liquidus temperature for alloy C M S X - 4 (1390 °C) was found to be 0.0436. The solubility of nitrogen at various pressures can be calculated as presented in Table 15. At the liquidus temperature of the alloy (1390 °C) , the solubility of nitrogen is found to be 8 ppm at a pressure of 3.5 x 10"5 atm. 121 Solubility of T i N : The dissolution of T i N in nickel based alloys is given by equation 21. The free energy for the reaction is given by [67]: AG" = 308780 - 114.35T Joules (45) The activity coefficients for T i and nitrogen were determined to be 0.727 and 0.323 respectively. The solubility constant K was calculated to be 1.88 x 10"4 (at the liquidus temperature). P N 2 (atm) [N] ppm 5 x 10 - 4 30 8 x 10"5 12 3.5 x 10"5 8 7 x 10"6 3.5 5 x IO"7 1 Table 15. Relation between the equilibrium nitrogen contents required for the formation of T i N and P N 2 for alloy C M S X - 4 (at 1390°C) . T ( ° C ) Superheat (°C) [N] ppm 1390 0 8 1440 50 15 1490 100 28 1540 150 51 Table 16. Equilibrium nitrogen content required for T i N formation in alloy C M S X - 4 above T i i q ( 1 3 9 0 ° C ) . 122 Table 16 presents the equilibrium nitrogen contents required for the formation of T i N in alloy C M S X - 4 , above the liquidus temperature (1390 °C) , at a pressure of 3.5 x 10 ~ 5 Torr. It can be seen that the solubility of T i N in alloy C M S X - 4 increases with an increase in the superheat. 123 A p p e n d i x : 7 Apart from the T i N inclusions, carbides, carbo-nitrides, delta and laves phases have been observed. The carbides are niobium carbides which are invariable cubic in structure. They are grayish or bluish gray in color under an optical microscope. Figure 67 presents an S E M picture of some niobium-carbides found in the DS 718 sample. The carbo-nitrides mostly have an inner-most alumina core surrounded by a T i N phase, which are together enveloped by a grayish niobium carbide phase. Note that T i N particles are precipitated first and the carbides are precipitated later during solidification (below the carbide solution temperature of 1280 °C) . These are large precipitates > 10 microns in size and are near the bottom of the directionally solidified ingot (due to phase precipitation equilibrium). Figure 68 presents a typical carbo-nitride found in the D S cast 718 samples. The laves phase and the needle shaped delta phases are niobium rich phases. They are the NiaNb phases distributed through out the DS cast ingot. S 3 2 9 i 5 Figure 67. S E M picture of Niobium Carbides. 124 Figure 68. A typical carbonitride (TiN with an alumina core and surrounded by Nb carbide phase) (a) 125 032915 20KV Xl'.kbk' kt'.tHm (b) Figure 69 (a) Delta phase, (b) needle shaped delta phase surrounding a T i N inclusion (c) and (d) T i N inclusion enveloped by Nb rich Laves phases. 126 Mostly the niobium rich phases are found enveloping the T i N particles. Figures 69 (a), (b), (c), and (d) show the needle shaped delta (NiaNb), niobium rich Laves phases enveloping titanium nitrides. 127 Appendix : 8 The very top of the D S cast IN718 ingots invariably have a chromium rich deposit. Figure 70 shows a S E M picture of the very top of the sample. A combination of high temperature and low pressure is effective in vaporizing the chromium from the ingot and causing a chromium rich deposit on the walls of the crucible above the ingot and also on the top of the ingot. Figure 71 shows a large T i N particle amidst chromium globules found on the upper outer surface of a DS castfN718 ingot. Figure 70. Chromium rich deposit on the top Figure 71. A T i N particle amidst chromium of the ingot. globules on the upper outer surface of the DS cast ingot. 128 

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