Interfacial microstructure and mechanical properties of diffusion-bonded joints of titanium TC4 (Ti-6Al-4V) and Kovar (Fe-29Ni-17Co) alloys

Diffusion bonding is a near net shape forming process that can join dissimilar materials through atomic diffusion under a high pressure at a high temperature.Titanium alloy TC4(Ti-6 Al-4 V)and 4 J29 Kovar alloy(Fe-29 Ni-17 Co)were diffusely bonded by a vacuum hot-press sintering process in the temperature range of 700-850°C and bonding time of 120 min,under a pressure of 34.66 MPa.Interfacial microstructures and intermetallic compounds of the diffusion-bonded joints were characterized by optical microscopy,scanning electron microscopy,X-ray diffraction(XRD)and energy dispersive spectroscopy(EDS).The elemental diffusion across the interface was revealed by electron probe microanalysis.Mechanical properties of joints were investigated by micro Vickers hardness and tensile strength.Results of EDS and XRD indicated that(Fe,Co,Ni)-Ti,TiNi,Ti_2Ni,TiNi_2,Fe_2 Ti,Ti_(17) Mn_3 and Al_6 Ti_(19) were formed at the interface.When the bonding temperature was raised from 700 to 850°C,the voids of interface were reduced and intermetallic layers were widened.Maximum tensile strength of joints at 53.5 MPa was recorded by the sintering process at 850°C for 120 min.Fracture surface of the joint indicated brittle nature,and failure took place through interface of intermetallic compounds.Based on the mechanical properties and microstructure of the diffusion-bonded joints,diffusion mechanisms between Ti-6 Al-4 Vtitanium and Fe-29 Ni-17 Co Kovar alloys were analyzed in terms of elemental diffusion,nucleation and growth of grains,plastic deformation and formation of intermetallic compounds near the interface.     

Effect of long-term room-temperature storage on the structure and properties of glassy melt-spun Mg-Al-Ca alloys

The effects of 48 months storage at room temperature on the structure and properties of glassy meltspun Mg-Al-Ca alloys has been determined. The results show that while Knoop hardness does not differ significantly from values obtained shortly after melt spinning, x-ray diffraction (XRD) shows considerably stronger signals from αMg present in the glassy matrix, detected only for the most dilute glass forming composition some months after melt spinning. High-resolution electron microscopy (HREM) studies show the presence of nanocrystallites in the glassy matrix with interplanar spacings matching αMg, and energy-dispersive spectroscopy (EDS) indicates the presence of Al and Ca within these nanocrystallites. Differential scanning calorimetry (DSC) studies at 20, 40, and 80 K/min show evidence of two-stage crystallization with the first stage starting at ∼120 deg with a peak at ∼150 °C and the second stage peaking at ∼250 °C. The significance of these results is discussed.     

Effect of reaction products on mechanical properties of diffusion bonded of titanium to 304 stainless steel with Cu interlayer joints

In present study, pure copper was used as an interlayer of the diffusion bonded joints between commercially pure titanium and 304 stainless steel. The process was carried out at 900°C for 30–150 minutes in steps of 30 minutes under 3MPa uniaxial load in vacuum. Microstructures of the bonded assemblies were observed in optical and scanning electron microscopes. The study exhibits the presence of different reaction layers in the diffusion zone and their chemical compositions were determined by energy dispersive spectroscopy. Formation of reaction products in diffusion interfaces were confirmed by xray diffraction technique. The maximum tensile strength of ∼322MPa (∼101% of Ti) and shear strength of ∼254MPa (∼86% of Ti) along with 8.4 % ductility were obtained for the couple bonded for 60 minutes and due to the rise in bonding time up to 90 minutes, bond strength drops marginally. With a further increase in joining time, the bond strength drops gradually due to the increase in the width of Fe-Ti base reaction products. Observation of fracture surfaces in SEM using EDS indicated that fracture takes place through the SS-Cu interface for all bonded samples.     

Embrittlement of titanium alloys by long time,high temperature exposure

α stabilized titanium alloys are known to exhibit embrittlement after long-time exposures above ∼800°F. The time-temperature dependency of this embrittlement phenomenon in the Ti-6Al-2Sn-4Zr-2Mo and Ti-5Al-6Sn-2Zr-lMo-0.25Si alloys was observed using a substandard fracture mechanics test. Room temperature slow-bend tests of fatigue precracked Charpy specimens were used to monitor toughness degradation after unstressed thermal exposures in the temperature range of 800° to 1100°F for times to 5000 hr. The activation energy for the embrittlement process was found to be ∼25 to 28 kcal per g mole, which approximates that for diffusion of aluminum or tin in α-Ti. The embrittlement is attributed to the Ti₃X (X = Al, Sn) phase with the rate controlling step that of diffusion controlled growth of the Ti₃X phase domains. The embrittlement process is reversible by heat treatment at temperatures above the α + Ti₃X two phase region.     

Growth of semicoherent TiFe nanoparticles in β-Ti matrix in the Ti73Fe27 rapidly quenched ribbon

A thin Ti₇₃Fe₂₇ ribbon was prepared by rapid quenching from the melt. The as-quenched ribbon was in a metastable condition with a small amount of nanoparticles of TiFe, of a size of 138±31 nm, embedded in the β-Ti matrix. The β-Ti matrix was supersaturated with Fe, and the fraction of the matrix was higher than that in the equilibrium state. High-resolution imaging of the interface of the TiFe/β-Ti showed that a periodic array of dislocations were present in the interface to accommodate the lattice mismatch. The spacing between the dislocations in the as-quenched specimen was 5.0±0.5 nm. When the ribbon was heated to 700 °C, growth of the TiFe nanoparticles to a size of 228±37 nm took place in the β-Ti matrix. The amount of β-Ti was reduced, as well as the Fe content in β-Ti. The interface between the TiFe and β-Ti remained semicoherent, except that the spacing between the interfacial dislocations was reduced to 3.5±0.6 nm. The persistence of the semicoherent interface was ascribed to the same crystal structure and close lattice parameters shared by TiFe and β-Ti. The growth kinetics of the TiFe nanoparticles during heating was examined based on the modified theory of isothermal heating. It can be considered to be controlled by the diffusion of Fe atoms in the β-Ti matrix to the TiFe phase. Prolonged heating of the ribbon below the eutectoid temperature led to partial transformation of β-Ti to α-Ti.     

Iron intermetallic phases in the Al corner of the Al-Si-Fe system

The iron intermetallics observed in six dilute Al-Si-Fe alloys were studied using thermal analysis, optical microscopy, and image, scanning electron microscopy/energy dispersive X-ray, and electron probe microanalysis/wavelength dispersive spectroscopy (EPMA/WDS) analyses. The alloys were solidified in two different molds, a preheated graphite mold (600 °C) and a cylindrical metallic mold (at room temperature), to obtain slow (∼0.2 °C/s) and rapid (∼15 °C/s) cooling rates. The results show that the volume fraction of iron intermetallics obtained increases with the increase in the amount of Fe and Si added, as well as with the decrease in cooling rate. The low cooling rate produces larger-sized intermetallics, whereas the high cooling rate results in a higher density of intermetallics. Iron addition alone is more effective than either Si or Fe+Si additions in producing intermetallics. The alloy composition and cooling rate control the stability of the intermetallic phases: binary Al-Fe phases predominate at low cooling rates and a high Fe:Si ratio; the β-Al₅FeSi phase is dominant at a high Si content and low cooling rate; the α-iron intermetallics (e.g., α-Al₈Fe₂Si) exist between these two; while Si-rich ternary phases such as the δ-iron Al₄FeSi₂ intermetallic are stabilized at high cooling rates and Si contents of 0.9 wt pct and higher. Calculations of the solidification paths representing segregations of Fe and Si to the liquid using the Scheil equation did not conform to the actual solidification paths, due to the fact that solid diffusion is not taken into account in the equation. The theoretical models of Brody and Flemings^[44] and Clyne and Kurz^[45] also fail to explain the observed departure from the Scheil behavior, because these models give less weight to the effect of solid back-diffusion. An adjusted 500 °C metastable isothermal section of the Al-Si-Fe phase diagram has been proposed (in place of the equilibrium one), which correctly predicts the intermetallic phases that occur in this part of the system at low cooling rates (∼0.2 °C/s).     

The Al-Cu-Fe phase diagram: 0 to 25 At. pct Fe and 50

Isothermal sections of the Al-Cu-Fe equilibrium phase diagram at temperatures from 680 °C to 800 °C were determined in the region with 50 to 75 at. pct Al and 0 to 25 at. pct Fe using scanning electron microscopy/energy dispersive spectroscopy (SEM/EDS) techniques. This re- gion includes the face-centered icosahedral phase (Ψ-Al₆Cu₂Fe) which has unprecedented struc- tural perfection and no apparent phason strain. The icosahedral phase has equilibrium phase fields with four distinct phases at 700 °C and 720 °C (β-Al(Fe, Cu), λ-Al₁₃Fe₄, ω-Al₇Cu₂Fe, and liquid) and three phases at 680 °C(β, ω, and λ) and 800 °C (β, λ, and liquid). The B2 ordered β phase has considerably greater solubility for Cu than previously reported, extending from AlFe to ∼Al₅₀Fe₅Cu₄₅. The equilibrium range of composition for the icosahedral phase at these temperatures was determined, and a liquidus projection is proposed.     

Diffusion in the titanium-nickel system: I. occurrence and growth of the various intermetallic compounds

Various types of diffusion couples have been prepared and the growth of the layers, formed in these couples, has been studied in the temperature range between 550 and 940°C. All phases which, on the basis of the equilibrium diagram, could be expected under the circumstances, were found in the diffusion zone. The layer growth of β-Ti(Ni), Ti₂Ni and TiNi in all types of couples obeyed the parabolic law; the TiNi layer was the only exception. It is supposed that in this case grain boundary diffusion is responsible for the observed abnormal growth. The shape of the β-Ti/β-Ti + Ti₂Ni boundary in the equilibrium diagram was constructed using data, both from recorded penetration curves and from equilibrated two-phase alloys. Excellent agreement between the results of these two methods was obtained. The eutectoid composition at 77Q°C has been placed at about 7 at. pct Ni, in contrast with earlier results. Furthermore, the composition range of TiNi has been determined from penetration curves. This paper is based on a Thesis submitted by G. F. BASTIN in fulfillment or requirements for the degree of Doctor in Technological Sciences.     

Time-temperature-precipitation characteristics of high-nitrogen austenitic Fe−18Cr−18Mn−2Mo−0.9N steel

The time-temperature-precipitation (TTP) and corresponding mechanical properties in high-nitrogen austenitic Fe−18Cr−18Mn−2Mo−0.9N steel (all in weight percent) were investigated using electron microscopy and ambient tensile testing. The precipitation reactions can be categorized into three stages: (1) high-temperature region (above 950°C)—mainly coarse grain-boundary (intergranular) Cr₂N; (2) nose-temperature region—integranular Cr₂N→cellular Cr₂N→intragranular Cr₂N+ sigma (σ); and (3) low-temperature region (below 750°C)—intergranular Cr₂N→cellular Cr₂N→ intragranular Cr₂N+σ+chi(χ)+M₇C₃ carbide. After cellular Cr₂N precipitation became dominant above 800°C, yield and tensile strength gradually decreased, whereas elongation abruptly deteriorated with aging time. On the contrary, prolonged aging in the low-temperature regime increased tensile strength, caused by the precipitation of fine χ and M₇C₃ within grains. Based on the analyses of selected area diffraction (SAD) patterns, the crystallographic features of the second phases were analyzed.     

Elevated temperature deformation behavior of a dispersion-strengthened Al-Fe,V,Si alloy

The deformation behavior of a rapidly solidified, dispersion-strengthened Al alloy containing 11.7 pct Fe, 1.2 pct V, and 2.4 pct Si was studied at test temperatures up to 450 °C using constantstress creep and constnt strain-rate tensile tests. Apparent stress exponents (n) up to ∼24 and an activation energy of 360 kJ/mol were obtained with the standard Arrhenius type power-law creep equation, which also suggested a change in behavior at ∼300 °C. Substructure-invariant and dislocation/dispersoid interaction models were found to be inadequate for explaining the behavior. When the data were replotted as vs σ, two regimes were found between 350 °C and 450 °C. A model with a pseudothreshold stress (σ _Th′ ) for the higher stress regime resulted inn ∼3, indicating solute drag in this regime. Transmission electron microscopy (TEM) showed departureside pinning of dislocations at higher stresses. In the lower stress regime, TEM showed dislocation subgrain structures. Here, the model resulted in a stress exponent of ∼4.5 indicating the dislocation climb mechanism. At temperatures below ∼300 °C, a single regime was found along with lower activation energies and a stress dependence of ∼3. Dislocation pipe diffusion is proposed to explain the lower activation energy. The origin ofσ _Th′ has been tied to dislocation generation at the grain boundaries.     

Effect of overaging on mechanical properties and stress corrosion cracking of HY-180M steel

The tensile properties, fracture toughness and stress corrosion cracking (SCC) behavior of HY-180 M steel at 22 °C were studied after final 5 h overaging treatments >510 ≤650 °C. SCC tests were conducted for 1000 h with compact tension specimens in aqueous 3.5 pct NaCl solutions at a noble (anodic) potential of −0.28 V_SHE ( −0.48 V_Ag/AgC1) and a cathodic protection potential of −0.80 V_SHE (−1.0 V_Ag/AgC1). The SCC resistance improved at aging temperatures >565 °C, the most significant improvement being at −0.80 V_She, especially after 650 ° aging whereK _ISCC was raised to at least 110 MPa · m^1/2. However, this was at the expense of mechanical properties. Provided low crack propagation rates of ∼3 X 10⁻¹¹ m/s at −0.80V _SHEmay be tolerated, the best compromise between strength, toughness, and SCC resistance was obtained after 594 °C aging. Under these conditions, stress intensities as high as ∼ 110 MPa · m^1/2 can be used, with a yield strength of ∼ 1150 MPa and fracture toughness of ∼ 170 MPa · m^1/2. The retained austenite content after aging increased with aging temperature up to 25 pct by vol at 650 °C. It appeared to correlate with improved SCC resistance, but other microstructural effects associated with aging may be involved. Formerly Research Associate with theDepartment of Metallurgical Engineering , University of BritishColumbia     

Effect of heat treatment on the microstructure,tensile properties,and fracture behavior of permanent mold Al-10 wt pct Si-0.6 wt pct Mg/SiC/10p composite castings

The present study was undertaken to investigate the effect of solution treatment (in the temperature range 520 °C to 550 °C) and artificial aging (in the temperature range 140 °C to 180 °C) on the variation in the microstructure, tensile properties, and fracture mechanisms of Al-10 wt pct Si-0.6 wt pct Mg/SiC/10_p composite castings. In the as-cast condition, the SiC particles are observed to act as nucleation sites for the eutectic Si particles. Increasing the solution temperature results in faster homogenization of the microstructure. Effect of solution temperature on tensile properties is evident only during the first 4 hours, after which hardly any difference is observed on increasing the solution temperature from 520 °C to 550 °C. The tensile properties vary significantly with aging time and temperature, with typical yield strength (YS), ultimate tensile strength (UTS), and percent elongation (EL) values of ∼300 MPa, ∼330 MPa, and ∼1.4 pct in the underaged condition, ∼330 MPa, ∼360 MPa, and ∼0.65 pct in the peakaged condition, and ∼323 MPa, ∼330 MPa, and ∼0.8 pct in the overaged condition. Prolonged solution treatment at 550 °C for 24 hours results in a slight improvement in the ductility of the aged test bars. The fracture surfaces exhibit a dimple morphology and cleavage of the SiC particles, the extent of SiC cracking increasing with increasing tensile strength and reaching a maximum in the overaged condition. Microvoids act as nucleation sites for the formation of secondary cracks that promote severe cracking of the SiC particles. A detailed discussion of the fracture mechanism is given.     

Kinetics of the sulfidation of chalcopyrite with gaseous sulfur

The sulfidation of chalcopyrite with gaseous sulfur in the temperature range of 325 °C to 400 °C occurs with the formation of covellite and pyrite as the final products. The rate of sulfidation depends strongly on the temperature, with nearly complete conversion in less than 30 minutes at 400 °C. Microscopic analysis of partially and completely reacted particles showed that the sulfidation proceeded topochemically, with a shrinking core of unreacted chalcopyrite surrounded by successive layers of FeS₂ and CuS. The experimental data exhibited an induction period at the beginning of the reaction. An electrochemical mechanism is proposed for the sulfidation reaction, which involves simultaneous diffusion of Cu⁻ and electrons through the product layers. The rate data showed that the fraction reacted is well represented by a shrinking-core model controlled by the reaction occurring at the chalcopyrite-pyrite interface, resulting in the conversion-vs-time relationship 1−(1−X)^1/3=k(t−t _ind). An activation energy of 98.4 kJ/mol was determined for the temperature range of 325 °C to 400 °C.     

Low-Temperature Processing of ZrB2-ZrC Composites by Reactive Hot Pressing

Dense ZrB₂-ZrC and ZrB₂-ZrC _x∼0.67 composites have been produced by reactive hot pressing (RHP) of stoichiometric and nonstoichiometric mixtures of Zr and B₄C powders at 40 MPa and temperatures up to 1600 °C for 30 minutes. The role of Ni addition on reaction kinetics and densification of the composites has been studied. Composites of ∼97 pct relative density (RD) have been produced with the stoichiometric mixture at 1600 °C, while the composite with ∼99 pct RD has been obtained with excess Zr at 1200 °C, suggesting the formation of carbon deficient ZrC _x that significantly aids densification by plastic flow and vacancy diffusion mechanism. Stoichiometric and nonstoichiometric composites have a hardness of ∼20 GPa. The grain sizes of ZrB₂ and ZrC _x∼0.67 are ∼0.6 and 0.4 μm, respectively, which are finer than those reported in the literature. This article is based on a presentation given in the symposium entitled “Materials Behavior: Far from Equilibrium” as part of the Golden Jubilee Celebration of Bhabha Atomic Research Centre, which occurred during December 15–16, 2006 in Mumbai, India.     

Diffusional reactions during processing of TIMETAL 21S/Al2O3 composites

The stability of an Al₂O₃ reinforcement in TIMETAL 21S has been investigated by annealing diffusion couples and consolidated fiber composites at 1100 °C, 900 °C, and 750 °C. Diffusion couple studies indicate that γ-TiAl, α ₂-Ti₃Al, and α-Ti(Al,O) phases can form upon annealing above the β transus of TIMETAL 21S, but γ-TiAl, α ₂-Ti₃Al, and a ternary T phase form during annealing below the β transus. The phases developed during diffusional interaction define a diffusion path between TIMETAL 21S and Al₂O₃. A coating of Nb, Mo, or Ta between TIMETAL 21S and Al₂O₃ acts as a diffusion barrier, but the coatings can diffuse into TIMETAL 21S at high temperature. In agreement with a kinetics analysis, a 2-μm-thick interface coating of Nb, Mo, or Ta in the TIMETAL 21S/Al₂O₃ composite can prevent the reaction during processing (2 hours at 850 °C or 900 °C) with no detectable diffusion into the matrix. If there are imperfections such as pinholes or cracks present in the diffusion barrier, the reaction quickly starts at the interface and does not remain confined at the imperfection; rather, it progresses along the interface. The mechanism for progressive development of interface reaction at a discontinuity in the diffusion barrier has been proposed. The analysis of the diffusional interface reactions in this work has identified some of the governing design concepts for development of robust high-temperature titanium-based composites.     

Diffusion Bonding of 17-4 Precipitation Hardening Stainless Steel to Ti Alloy With and Without Ni Alloy Interlayer: Interface Microstructure and Mechanical Properties

In the present study, the diffusion bonding of 17-4 precipitation hardening stainless steel to Ti alloy with and without nickel alloy as intermediate material was carried out in the temperature range of 1073 K to 1223 K (800 °C to 950 °C) in steps of 298 K (25 °C) for 60 minutes in vacuum. The effects of bonding temperature on interfaces microstructures of bonded joint were analyzed by light optical and scanning electron microscopy. In the case of directly bonded stainless steel and titanium alloy, the layerwise α-Fe + χ, χ, FeTi + λ, FeTi + β-Ti phase, and phase mixture were observed at the bond interface. However, when nickel alloy was used as an interlayer, the interfaces indicate that Ni₃Ti, NiTi, and NiTi₂ are formed at the nickel alloy-titanium alloy interface and the PHSS-nickel alloy interface is free from intermetallics up to 1148 K (875 °C) and above this temperature, intermetallics were formed. The irregular-shaped particles of Fe₅Cr₃₅Ni₄₀Ti₁₅ have been observed within the Ni₃Ti intermetallic layer. The joint tensile and shear strength were measured; a maximum tensile strength of ~477 MPa and shear strength of ~356.9 MPa along with ~4.2 pct elongation were obtained for the direct bonded joint when processed at 1173 K (900 °C). However, when nickel base alloy was used as an interlayer in the same materials at the bonding temperature of 1148 K (875 °C), the bond tensile and shear strengths increase to ~523.6 and ~389.6 MPa, respectively, along with 6.2 pct elongation.     

Diffusion in the titanium-nickel system: II. calculations of chemical and intrinsic diffusion coefficients

Diffusion coefficients in the Ti-Ni system have been calculated by the aid of equations given by Sauer and Freise, and Wagner. Values for the TiNi (50 at. pct Ni) phase were found to be:D ^u (cm²/s) = 0.0020 exp - 142,000/R for the temperature range between 650 and 940°C. The heat of activation, expressed in J/mol, has an accuracy of ±6000. For the β-Ti(Ni) phase containing 6 at. pct Ni the temperature dependence of the diffusion coefficient is expressed by:D ^u (cm²/s) = 0.0688 exp - 141,000/RT. The uncertainty in the energy of activation is ±12000 J/mol. No clear variation of the diffusion coefficient with concentration could be detected. It was found that Ni is by far the fastest moving component in β-Ti(Ni), Ti₂Ni and TiNi (at least in the composition range between 50 and 53 at. pct Ni). Values ofD _Ni/D _Ti have been calculated with an equation derived by van Loo. The significance of the calculated values is critically examined. By means of a practical example it is shown that the calculated ratio of the intrinsic diffusion coefficients can be extremely sensitive to slight variations in the position of the marker interface.Diffusion coefficients in the Ti-Ni system have been calculated by the aid of equations given by Sauer and Freise, and Wagner. Values for the TiNi (50 at. pct Ni) phase were found to be:D ^u (cm²/s) = 0.0020 exp - 142,000/R for the temperature range between 650 and 940°C. The heat of activation, expressed in J/mol, has an accuracy of ±6000. For the β-Ti(Ni) phase containing 6 at. pct Ni the temperature dependence of the diffusion coefficient is expressed by:D ^u (cm²/s) = 0.0688 exp - 141,000/RT. The uncertainty in the energy of activation is ±12000 J/mol. No clear variation of the diffusion coefficient with concentration could be detected. It was found that Ni is by far the fastest moving component in β-Ti(Ni), Ti₂Ni and TiNi (at least in the composition range between 50 and 53 at. pct Ni). Values ofD _Ni/D _Ti have been calculated with an equation derived by van Loo. The significance of the calculated values is critically examined. By means of a practical example it is shown that the calculated ratio of the intrinsic diffusion coefficients can be extremely sensitive to slight variations in the position of the marker interface. This paper is based on a Thesis submitted by G. F. BASTIN in fulfillment of requirements for the degree of Doctor in Technological Sciences.     

Transformation of Alumina Inclusions by Calcium Treatment

The objectives of this study were to investigate reactions of calcium with Al₂O₃ by different model experiments both on the laboratory and on the industrial scale. Experiments with solid Al₂O₃ and CaO were performed between 1350 °C and 1600 °C. Reaction rate constants were determined based on scanning electron microscopy/energy dispersive spectroscopy (SEM/EDS) observations of reaction products and weight measurements of the Al₂O₃ reacted via dissolution of the CaO bearing phases from the specimens after the annealing period. The results showed that the formation of calcium aluminate phases proceeded rapidly at temperatures greater than 1405 °C when a liquid calcium aluminate was formed. In the lowest temperature range (1350 °C–1405 °C), when the formation of liquid phase ceased, the reaction rate was several orders of magnitude lower. Industrial trials including Ca-alloy injection into steel, sampling and SEM/EDS analyses, as well as an inclusion rating in the samples show the concept of rapid transformation of the alumina inclusions with Ca treatment.     

The tantalum-palladium constitution diagram

The Ta-Pd system was investigated over the entire composition range by metallography, X-ray diffraction and electron microprobe analysis. The solubility limits of terminal and intermediate phases and solidus temperatures were determined. α-Ta dissolved ∼20 at. pct Pd at 2550°C and ∼10 at. pct Pd at 1000°C; α-Pd dissolves ∼22 at. pct Ta at 1730°C and ∼18 at. pct Ta at 1000°C. The presence of four intermediate phases a, (β-U type), α-TaPd (TiCu type), TaPd₂ (MoPt₂ type), and TaPd₃ (TiAl₃ type) was confirmed; they melt or decompose (α-TaPd) at about 2550, 1410; 1800, and 1770°C, respectively. In addition, an equiatomic high temperature phase, β-TaPd was found which melts at ∼1720°C and may be an extension of and isomorphous with the α-Pd solution. Seven three-phase reactions are described. Formerly with Massachusetts Institute of Technology