Effect of gamma-gamma prime mismatch,volume fraction gamma prime,and gamma prime morphology on elevated temperature properties of Ni, 20 Cr, 5.5 Mo,Ti, Al alloys

This study was designed to provide a critical test for the postulate that the mismatch between the lattices of austenite (γ) and the age-hardening gamma-prime (γ′) precipitates and the resultant coherency strains have a significant influence on the elevated temperature, particularly stress rupture, properties of a nickel-base superalloy. Two experimental alloys with a base analysis of Ni, 20 Cr, 5.5 Mo were designed with variable titanium and aluminum additions. To discern the effect of mismatch, an alloy without molybdenum was also experimented with. By manipulating the mismatch and volume fraction γ′ by heat treatment and chemistry, it was shown that a lower γ-γ′ mismatch indeed is beneficial to stress rupture life. Importance of volume fraction γ′ on this elevated temperature was also established.     

Coherency strengthening in Ni base alloys hardened by DO22 γ′ precipitates

The strengthening phase in nickel-base Alloy 718 is a metastable DO₂₂ precipitate termed γ′ which is based upon Ni₃Nb. Since the coherent γ′ particles produce a tetragonal distortion of the matrix, the specific variants of γ′ present can be controlled by the application of stress during aging. The resultant variation in strength is appreciable. A detailed comparison with likely strengthening mechanisms indicates that hardening is primarily due to the coherency strains arising from the γ′ precipitate. Implications with regard to design of alloys strengthened by the DO₂₂ phase are considered.     

MAX phases in titanium and aluminum alloys

The crystalline structures in the Ti-Al-C and Ti-Si-C systems are analyzed, and experiments are conducted with VT6 titanium alloy and eutectoidal Silumin (Al-12% Si) subjected to electroexplosive alloying and electron-beam treatment. Diffraction analysis reveals the formation of MAX phases (Ti₃SiC₂ and Ti₃AlC) in the modified layer of these alloys.     

Measurement and analysis of distribution coefficients in Fe-Ni alloys containing S and/or P: Part I.K Ni andK P

Fe-Ni alloys containing S and/or P were grown by a plane front solidification technique. The compositions of each element of these alloys were measured along the growth direction with an electron probe microanalyzer and were plotted as a function of fraction solidified. The distribution coefficient for each solute element was determined by an analytical procedure from the composition data as a function of fraction solidified. The distribution coefficient of Ni varies from 0.9 to 1.30. The distribution coefficient of P varies from 0.10 to 0.20 in a S-free Fe-Ni-P alloy, and from 0.25 to 5.80 in S-containing Fe-Ni-P alloys. This significant variation in the distribution coefficient of P in Fe-Ni-S-P alloys is attributed to the interactions occurring between S and P in the liquid during solidification. The distribution coefficient of Ni increases with the S and P contents of the liquid while the distribution coefficient of P increases with the S content and decreases with the P content of the liquid. Mathematical expressions have been formulated to relate the distribution coefficients of Ni and P to the S and P contents in the liquid.     

The nucleation of aluminum grains in alloys of aluminum with titanium and boron

The microstructure of a ternary alloy, Al-5 wt pct Ti, 1 wt pct B, has been examined by optical and electron transmission microscopy, by selected area diffraction, and electron probe microscopy, by selected area diffraction, and electron probe microanalysis. Particles of Al₃Ti are found at the center of grains and there exist preferred epitaxial orientations between this compound and the surrounding aluminum. Particles containing titanium and boron occur at aluminum grain boundaries and have no preferred configurations with respect to the aluminum or to one another. It is concluded that the active heterogeneous nuclei are therefore Al₃Ti and that particles of TiB₂, AlB₂, or a ternary compound are not active in this alloy. Grain size measurements in binary Al-Ti alloys suggest that particles of a nucleating phase must be present at concentrations as low as 0.01 wt pct Ti, and it is suggested that these could be Al₃Ti if the existing binary phase diagram Al-Ti is in error.     

Combined matrix/boundary precipitation strengthening in creep of Fe-15 Cr-25 Ni alloys

High temperature creep behavior of carbide precipitation strengthened Fe-15 Cr-25 Ni alloys with different carbon content have been investigated. Grain boundary carbides obstruct dislocation annihilation at the grain boundary and, therefore, increase the dislocation density near the grain boundary. This gives rise to formation of a hard grain boundary region and significantly increase creep resistance of the alloy. The grain boundary precipitation strengthening and combined matrix/boundary strengthening are modeled following the concept of hard-soft composite structure, and a unified creep equation is derived by taking account of back stress from intergranular carbide particles, “boundary obstacle stress”. The models and analysis show that grain boundary precipitation strengthening is predominant for soft matrix but decreases with the increase of matrix strength, indicating the existence of coupled matrix/boundary strengthening.     

The growth of gamma prime precipitates in aged Ni−Ti alloys

The kinetics of growth of the γ′ precipitate in a Ni-8.74 wt pct Ti alloy were studied by magnetic analysis and transmission electron microscopy. The variation of the titanium content of the nickel-rich matrix as a function of aging time was studied by measuring the ferromagnetic Curie temperature of alloys aged at 692°, 593°, and 525°C. The kinetics of this process accurately obeyed the predictions of the Lifshitz-Wagner theory of diffusion controlled coarsening after relatively short aging times at all aging temperatures. Dark-field transmission electron microscopy was used to measure the particle-size distributions and the average particle sizes of samples aged for various times at 692°C. The kinetics of particle growth also obeyed the time law predicted by the Lifshitz-Wagner theory within the limits of experimental error. Additional analysis of the data provided a value of approximately 21 erg per sq cm for the interfacial free energy of the γ′-matrix interface, and a value for the diffusion coefficient of titanium in nickel which is in very good agreement with an independently determined value. The distribution of γ′ particle sizes was found to be significantly broader than the theoretical distribution of the Lifshitz-Wagner theory. It is suggested that this is due to the relatively large lattice parameter mismatch between γ′ and the Ni−Ti matrix. The results and conclusions of this study are critically compared with those of other investigations. A. J. ARDELL, formerly Assistant Professor, California Institute of Technology     

Nitrogen solubility and aluminum nitride precipitation in liquid iron alloys containing nickel and aluminum

The solubility of nitrogen in liquid iron-base Fe-Ni-Al alloys has been measured up to the solubility limit for formation of aluminum nitride using the Sieverts’ method. Measurements were conducted over the temperature range from 1843 to 2023 K and aluminum concentration range from 1.5 to 3.0 wt pct Al. The effect of nickel additions was determined at 2, 5 and 10 wt pct Ni. The cross interaction parameter describing the effect of nickel and aluminum on the activity coefficient of nitrogen in iron was determined. The first and second order effects of nickel on the activity coefficient of aluminum also were determined. The solubility product of aluminum nitride increases with increasing aluminum content and increasing temperature. Addition of nickel decreases the solubility products of aluminum nitride in lower aluminum content alloys. However, the effect of the cross interaction terme _Al ^NiAl becomes significant with increasing aluminum content and compensates for the effects of the first and second order nickel-nitrogen and nickelaluminum interaction terms. Therefore the effect of nickel additions show little effect on the solubility products of aluminum nitride in higher aluminum alloys.     

Mechanical properties of Ir-Nb alloys containing Ni and Al

Two alloys made by adding 5 or 10 at. pct, respectively, of Ni-18.9 at. pct Al to an Ir-15 at. pct Nb alloy were investigated. The microstructure and compressive strength at temperatures between room temperature and 1800 °C were investigated to evaluate the potential of these alloys for ultra-high-temperature use. Their microstructural evolution indicated that the two alloys formed fcc and L1₂-Ir₃Nb two-phase structures. The fcc and L1₂ two-phase structures were examined by transmission electron microscopy (TEM) and scanning electron microscopy (SEM). The 0.2 pct flow stresses were above 1000 MPa at temperatures up to 1200 °C, about 150 MPa at 1500 °C, and over 100 MPa at 1800 °C. The strength of the quaternary Ir-base alloys at 1200 °C was even higher than that of Ir-base binary and ternary alloys. And the strength of quaternary Ir-Nb-Ni-Al was equivalent to that of the Ir-15 at. pct Nb binary alloy at 1800 °C. The compressive ductility of quaternary (around 20 pct) was improved drastically compared with that of the Ir-base binary alloy (lower than 10 pct) and the ternary Ir-base alloys (about 11 pct). An excellent balance of high-temperature strength and ductility was obtained in the alloy with 10 at. pct Ni-18.9 at. pct Al. The effect of Ni and Al on the strength of the Ir-Nb binary alloy is discussed.     

Thermodynamics of oxygen solutions in nickel melts containing aluminum and titanium

Thermodynamic analysis of oxygen solutions in nickel melt shows that, as aluminum and titanium are added to the melt, the solubility of oxygen decreases. However, after reaching 0.205% Al and 0.565% Ti, the oxygen concentration in the melt begins to rise with increase in the Al and Ti content. The minimum oxygen concentrations in the reduction of nickel melt by aluminum (1.44 × 10^–4% O) and titanium (2.98 × 10^–4% O) are determined. On that basis, we may propose the optimal approach to alloying nickel melts with aluminum and titanium. First, the melt is reduced by adding sufficient aluminum to minimize the oxygen concentration in the melt (~0.2% Al). Then the oxide formed is removed, so as to prevent repeated oxidation of the melt. Finally, the melt is alloyed with aluminum and titanium to obtain the required alloy composition.     

Diffusivity and solubility of hydrogen as a function of composition in Fe-Ni alloys

An electrochemical method has been used to determine the permeability,P, diffusion coefficient,D, and solubility,c, of hydrogen in alloys of the Fe-Ni system. The heats of activation for diffusion and the heats of solution have been derived.D falls from ≃10⁻⁴ sq cm per sec for pure iron to ≃10⁻¹⁰ sq cm per sec for 40 wt pct of Ni in the alloy. Thereafter it rises slightly to that for pure nickel,c rises by about 10³ between pure iron and 40 wt pct Ni, then remains constant up to pure nickel. The resultantP doubles at 5 wt pct Ni and then falls by 10³ times up to 40 wt pct Ni, afterwards rising slightly to that for pure nickel. Between 0 and 40 wt pct Ni the dominant factor in controlling the value ofP is the fall of the mole fraction of the α phase in the alloy. This hypothesis gives a reasonable quantitative calculation of theP-composition relation. Between 40 and 100 wt pct, the crystallographic phase is allγ and the major effect is the bonding of hydrogen in the alloy, the small changes noted being reasonably calculable. The negligible change of solubility in this region reflects the negligible change ind character of the alloy from 40 to 100 wt pct Ni. The hydrogen permeability of Fe-Ni (5 wt pct) is greater than that of palladium atT > 200°C. The corrosion rate and hydrogen permeability (hence, susceptibility to hydrogen embrittlement) pass through a minimum at about 50 wt pct Ni. A remarkable parallelism exists between corrosion rate and hydrogen permeation in Fe-Ni alloys. An interpretation is suggested. Formerly with the University of Pennsylvania Formerly with the University of Pennsylvania Work carried out by P. K. SUBRAMANYAN in partial fulfillment of the requirements for the degree of Doctor of Philosophy, University of Pennsylvania, 1970.     

Substructure of ausformed martensite in Fe-Ni and Fe-Ni-C alloys

The martensite substructure after ausforming has been studied for two different martensite morphologies: partially twinned, lenticular martensite (Fe-33 pct Ni, M_s =-105‡C) and completely twinned “thin plate” martensite (Fe-31 pct Ni-0.23 pct C, M_s = -170‡C), and in both cases ausforming produces a dislocation cell structure in the austenite which is inherited, without modification, by the martensite. In the Fe-Ni alloy, the dislocation cell structure is found in both the twinned (near the midrib) and untwinned (near the interface) regions, the latter also containing a regular dislocation network generated by the transformation itself and which is unaltered by the austenite dislocation cell structure. Similarly, in the Fe-Ni-C alloy, the transformation twins are unimpeded by the prior cell structure. These observations show that carbide precipitation during ausforming is not necessarily required to pin the austenite cell structure and that the martensite-austenite interface, backed by either twins or dislocations, does not exhibit a ”sweeping” effect. Although the martensite transformation twins are not inhibited by the ausforming cell structure, they do undergo a refinement with increased ausforming, and it is indicated that the transformation twin width in martensite depends on the austenite hardness. However, the relative twin widths remain unchanged, as expected from the crystallographic theory. T. MAKI, Formerly with the University of Illinois     

Thermodynamics of titanium and nitrogen in an Fe-Ni melt

The equilibrium solubility of titanium and nitrogen in Fe-Ni melts was measured in the presence of pure solid TiN under various nitrogen pressures in the temperature range of 1843 to 1923 K. The activity coefficients of titanium and nitrogen relative to a 1 mass pct standard state in liquid iron were calculated from the experimental results for Fe-Ni alloys of nickel contents up to 30 mass pct. Nickel decreases the activity coefficient of titanium, but it increases the activity coefficient of nitrogen in an Fe-Ni-Ti-N melt. Therefore, the effect of nickel on the solubility product of TiN is not significant. The first- and second-order interaction parameters of nickel on titanium (e _Ti ^Ni and r _Ti ^Ni , respectively) were determined to be −0.0115 and 0 at 1873 K, respectively. Similarly, the interaction parameters of nickel on nitrogen (e _N ^Ni and r _N ^Ni , respectively) were determined to be 0.012 and 0, respectively, at 1873 K. The temperature dependence of these interaction parameters was also determined.     

Creep resistance and high-temperature metallurgical stability of titanium alloys containing gallium

Tensile properties up to 1100°F and the creep resistance at 1000°F were correlated with composition for twelve complex developmental titanium alloys with additions of Al, Ga, Sn, Mo, Zr, and Si. Creep resistance for these alloys in the β heat-treated condition was found to be strongly dependent on the totalα stabilizer content and the silicon concentration. The creep activation energy for a Ti-4.5 Al-2 Sn-3 Zr-3 Ga-1 Mo-0.5 Si alloy, established over the 900° to 1100°F temperature range, was about 100 kcal per g-mole. This high creep activation energy is hypothesized to result from dispersion strengthening within theα matrix by the Ti₃ X (X = Al, Ga, Sn) phase and pinning of the interplatelet and priorβ grain boundaries by the Zr₅Si₃ phase. Both phases were identified by transmission electron microscopy in these respective locations. Metallurgical instability, as evidenced by decreased fracture toughness, is also shown to be relatable to the totalα stabilizer content. The activation energy for the embrittlement process is about 45 kcal per g-mole. which approximates that for interdiffusion of gallium inα titanium.     

The diffusivity of Ni in Fe-Ni and Fe-Ni-P martensites

The diffusivity of Ni in Fe-Ni and Fe-Ni-P martensite, , has been determined between 700 and 300 °C using electron microprobe (EMP) and scanning transmission electron microscope (STEM) techniques. Alloys of various bulk compositions (0 to 30 wt pct Ni, Fe) were homogenized in the single phase austenite (γ-fee) field and quenched to form martensite, α₂ (bcc). Appropriate alloys were tempered isothermally at 300 to 700 °C. The γ nucleated and grew in the parent α₂. The composition of the γ phase and the concentration gradients in the α₂ were measured with the EMP andJor STEM. In order to determine experimentally measured Ni concentration gradients were matched to Ni concentration gradients calculated by a simulation model. The calculated gradients were obtained by solving the appropriate form of Fick’s second law using the Crank-Nicholson numerical technique. The observed diffusivities varied with temperature. Above approximately 410 °C, while below 410°C, = (2.27 × 10⁻¹⁵) exp (− 10,600/RT) cm²/s. The effect of P is to increase the Fe-Ni diffusivities at any temperature by the factor (1 + 1.27C _p + 0.623C _p ² ) whereC _p is the amount of P (wt pct) in α₂. The discontinuous diffusion behavior of is attributable to the high dislocation density of the α₂. Above approximately 410 °C lattice diffusion is dominant while below 410 °C dislocation pipe diffusion is dominant. Formerly Research Assistant in the Department of Metallurgy and Materials Engineering, Lehigh University, Bethlehem, PA     

Substructure of ausformed martensite in Fe-Ni and Fe-Ni-C alloys

The martensite substructure after ausforming has been studied for two different martensite morphologies: partially twinned, lenticular martensite (Fe-33 pct Ni, M_s =-105?C) and completely twinned “thin plate” martensite (Fe-31 pct Ni-0.23 pct C, M_s = -170?C), and in both cases ausforming produces a dislocation cell structure in the austenite which is inherited, without modification, by the martensite. In the Fe-Ni alloy, the dislocation cell structure is found in both the twinned (near the midrib) and untwinned (near the interface) regions, the latter also containing a regular dislocation network generated by the transformation itself and which is unaltered by the austenite dislocation cell structure. Similarly, in the Fe-Ni-C alloy, the transformation twins are unimpeded by the prior cell structure. These observations show that carbide precipitation during ausforming is not necessarily required to pin the austenite cell structure and that the martensite-austenite interface, backed by either twins or dislocations, does not exhibit a ”sweeping” effect. Although the martensite transformation twins are not inhibited by the ausforming cell structure, they do undergo a refinement with increased ausforming, and it is indicated that the transformation twin width in martensite depends on the austenite hardness. However, the relative twin widths remain unchanged, as expected from the crystallographic theory.     

Heat flux transients at the casting/chill interface during solidification of aluminum base alloys

Heat flow at the metal/chill interface of bar-type castings of aluminum base alloys was modeled as a function of thermophysical properties of the chill material and its thickness. Experimental setup for casting square bars of Al-13.2 pct Si eutectic and Al-3 pet Cu-4.5 pct Si long freezing range alloys with chill at one end exposed to ambient conditions was fabricated. Experiments were carried out for different metal/chill combinations with and without coatings. The thermal history at nodal locations in the chill obtained during the experiments was used to estimate the interface heat flux by solving a one-dimensional Fourier heat conduction equation inversely. Using the data on transient heat flux q, the heat flow at the casting/chill interface was modeled in two steps: (1) The peak in the heat flux curve q_max was modeled as a power function of the ratio of the chill thickness d to its thermal diffusivity a, and (2) the factor (q/q_max) X α^0.05 was also modeled as a power function of the time after the solidification set in. The model was validated for Cu-10 pct Sn -2 pct Zn alloy chill and Al-13.2 pct Si and Al-3 pct Cu-4.5 pct Si as the casting alloys. The heat flux values estimated using the model were used as one of the boundary conditions for solidification simulation of the test casting. The experimental and simulated temperature distributions inside the casting were found to be in good agreement. Formerly Assistant Professor with Karnataka Regional Engineering College