Grain boundary pest of boron- doped Niin3Al at 1200 °C

A ductile Ni_77.4Al₂₂Zr_0.6B_0.2 alloy suffered severe intergranular embrittlement after air exposure at 1200 °C for 100 hours. However, the material survived after 1038 °C air exposure for 100 hours. Auger analysis showed enormous oxygen segregation on the grain boundaries in the 1200 °C, air exposed, boron-doped Ni₃Al, while the boron segregation remained unchanged. To elucidate this type of grain boundary damage, a modified fracture mechanism was proposed. Finally, anomalous grain growth was found in this alloy after 1200 °C air exposure, and an explanation for this phenomenon was suggested.     

Hydrogen embrittlement in ductile Ni3Al—Effects of hydrogen content,strain rate and pre-deformation

Hydrogen embrittlement has been studied in continuous cast sheet of an Ni₃Al alloy (Ni_77.83Al_21.73Zr_0.34B_0.1). When tensile tests were performed at the initial strain rate of 5.8 × 10⁻⁵ s⁻, the elongation decreased from 32.7% for no charging to 1.9% for 330 min of cathodic charging with 50 mA/cm² current, but the yield stress did not change. The fracture mode changed partially from dimple to intergranular and cleavage modes. At a faster strain rate of 5.8 × 10⁻³ s⁻¹, hydrogen embrittlement was less pronounced, but the yield stress increased with hydrogen content and multiple cracks were generated. Plastically pre-deformed (2–26% elongations) and subsequently charged specimens failed after yielding, which occurred at the final pre-deformation stress. Our results suggest that the grain boundary or the interior of the grain was not weakened by hydrogen, rather hydrogen-enhanced localized plasticity caused the loss of ductility in B-doped Ni₃Al alloy.     

Effect of Sulfur and Antimony on the Intergranular Fracture of Iron at Cathodic Potentials

Antimony, segregated to grain boundaries of iron, was found to be five times more effective than sulfur in promoting intergranular fracture of iron when tested in IN H₂SO₄ at cathodic potentials. A decrease in the ductility of iron accompanied the fracture mode change at increasing cathodic potentials. The effectiveness of antimony relative to sulfur was determined from straining electrode tests on iron and iron + 250 appm antimony alloys heat treated at 800 °C and 600 °C to produce different grain boundary chemical compositions. Grain boundary compositions were determined by Auger Electron Spectroscopy (AES). Similar grain boundary sulfur concentrations of 0.2 monolayers were observed by AES for the iron and iron + 250 appm antimony alloy after an anneal of 240 hours at 600 °C, while 0.08 monolayers of antimony was observed for the iron + 250 appm antimony alloy. These results suggest that sulfur and antimony do not compete for grain boundary sites.     

Improved Resistance to Laser Weld Heat-Affected Zone Microfissuring in a Newly Developed Superalloy HAYNES 282

Gleeble thermomechanical simulation and microstrucutural analyses of laser beam weldability of a newly developed precipitation-hardened nickel-base HAYNES alloy 282 were performed to better understand the fundamental cause of heat-affected zone (HAZ) cracking and how to prevent the cracking problem in the material. Submicron size intergranular M₅B₃ particles are identified for the first time in the present work by transmission electron microscopy, and were found to be the primary cause of HAZ grain boundary liquation cracking in the alloy. Complete dissolution of the liquating M₅B₃ particles by preweld heat treatment exacerbated rather than reduced susceptibility to cracking, which could be attributed to nonequilibrium intergranular segregation of boron atoms, liberated by the complete dissolution of the boride particles, during cooling from heat treatment temperature. Consequently, to reduce the HAZ cracking, a preweld heat treatment that reduces the volume fraction of the M₅B₃ particles while minimizing nonequilibrium grain boundary boron segregation is necessary, and this is possible by heat treating the alloy at 1353?K to 1373 K (1080?°C to 1100 °C). Further improvement in cracking resistance to produce crack-free welds is achieved by subjecting the alloy to thermomechanically induced grain refinement coupled with the preweld heat treatment at 1353 K (1080 °C). A Gleeble hot ductility test showed that formation of the crack-free welds is unexplainable by mere reduction in grain size without considering the effect of grain refinement on intergranular liquid produced by subsolidus liquation of the M₅B₃ borides.     

Electronic and structural studies of grain boundary strength and fracture in L12 ordered alloys—III. On the effect of stoichiometry

In order to clarify the effect of alloy stoichiometry on the grain boundary strength and fracture behavior of L1₂-type A₃B compounds, the mechanical properties of Ni₃Al, Ni₃Ga, Co₃Ti and Ni₃(Al_0.52Mn_0.48) compounds were investigated at room temperature. Also, the effect of interstitial atoms of boron and hydrogen and the interaction of its effect with alloy stoichiometry were observed. The result indicated that the magnitude of compositional dependence of alloy stoichiometry on the elongation, the ultimate tensile strength and fracture behavior varied as Ni₃Al > Ni₃Ga > Co₃Ti > Ni₃(Al_0.52Mn_0.48). It was also revealed that the effect of interstitial atoms on the grain boundary cohesive strength is independent on and then additive to the effect of alloy stoichiometry. The significance of the grain boundary structure and related electronic environment in the grain boundary region was, in the geometrical frame, emphasized to interpret the present behavior.     

High Cycle Fatigue of (Fe,Ni, Co)3V type alloys

High cycle fatigue tests in vacuum have been performed on ordered (Fe, Co, Ni)₃V alloys between 25 °C and 850 °C. Heat-to-heat variations in fatigue properties of a Co-16.5 wtpct Fe-25 pct alloy, LRO-1, appeared to be due to differing quantities of grain boundary precipitates. Modification of this alloy with 0.4 pct Ti, to produce an alloy designated LRO-23, reduced the density of grain boundary precipitates and increased ductility, resulting in superior fatigue strength at high temperatures. The fatigue lives of LRO-1 and LRO-23 decreased rapidly above 650 °C, and increased intergranular failure was noted. The fatigue resistance of a cobalt-free alloy, Fe-29 pct Ni-22 pct V-0.4 pct Ti (LRO-37), was examined at 25 °C, 400 °C, and 600 °C; there was little evidence for intergranular fracture at any of these temperatures. Fatigue behavior of the LRO alloys is compared to that of conventional high temperature alloys.     

Effects of boron doping on the grain-growth kinetics and mechanical properties of γ/γ′ nickel-aluminum alloys

The effects of 0.5 at. pct of boron doping on the microstructures and mechanical properties of γ/γ′ nickel-aluminum alloys have been investigated in the present study. A nickel-rich grain-boundary zone was observed in the boron-doped alloy after homogenization at 1100 °C and prolonged annealing at 1200 °C. Boron doping also caused remarkable improvements in toughness and tensile elongation and caused the fracture mode to change from completely intergranular to completely transgranular. The grain growth following recrystallization at 1200 °C was found to be retarded upon boron doping. A sudden increase in tensile elongation and a sudden drop in hardness were also observed upon prolonged heating during isothermal annealing at 1200 °C. The results are interpreted with reference to boron-nickel cosegregation at the grain boundaries.     

The influence of Mo and Co on the embrittlement of an Fe-Ni-Mn alloy

In the “as rolled” condition an Fe-6 Ni-5 Mn maraging type alloy was found to be brittle exhibiting intergranular fractures. The addition of 2.5 pct Mo and 5.0 pct Mo increased the impact toughness of the “as rolled” material and changed the mode of brittle fracture to transgranular cleavage. The addition of 9 pct Co embrittled the alloy. On aging Mo and Co raised the peak hardness of the base Fe-6 Ni-5 Mn alloy, however, aging led to rapid embrittlement. The base alloy and an alloy containing 2.5 pct Mo showed brittle intergranular fractures on aging. The addition of 5 pct Mo gave rise to brittle transgranular cleavage fractures on aging at 450°C, but at temperatures less than 450°C there was always up to 20 pct intergranular fracture present in brittle fractures. At temperatures greater than 475°C brittle intergranular failure occurred in the 5 pct Mo alloy due to a grain boundary film of M₆C and Fe₂Mo. This paper is based upon a thesis submitted by D. R. Squires in partial fulfilment for a higher degree of CNAA at Sheffield Polytechnic.     

Diffusion in the nickel-rich part of the Ni−Al system at 1000° to 1300°C; Ni3Al layer growth,diffusion coefficients,and interface concentrations

The formation of the Ni₃Al layer in NiAl (55 at. pct Ni)-pure Ni diffusion couples at temperatures above 1000°C has been found to be controlled almost completely by volume diffusion. At 1000°C and below, the relatively small grain size of the Ni₃Al compound in the layers caused such a large contribution from grain boundary diffusion, that the layer growth rates at 1000°C exceeded those at 1100°C and even those at 1200°C. In Ni₃Al (75at. pct Ni)-pure Ni diffusion couples the Ni₃Al compound rapidly converted into the solid solution of aluminum in nickel. Volume-diffusion coefficients calculated by the Boltzmann-Matano method yielded heats of activation of 55, 64, and 65 kcal·mol^?1 for NiAl, Ni₃Al and the solid solution of aluminum in nickel, respectively. In addition, eleven different types of diffusion couples were prepared from various Ni?Al alloys and annealed at 1000°C. Marker interface displacements and observations of porosity in these couples yielded a more detailed picture of the Kirkendall-effect than earlier work had done. The ratio of the intrinsic diffusion coefficients at the marker interface,D _NI/D _Al, is greater than one in the nickel-rich NiAl phase. For the Ni₃Al phase no statement can be made on the basis of this work. When the marker interface is located in the nickel solid solution,D _Ni/D _Al is smaller than one. The phase boundary concentrations in these couples did not show the expected deviation from the equilibrium concentrations in two-phase alloys; this finding is discussed with regard to the free-energycomposition diagram.     

Structure of As-cast aluminum matrix composite containing Ni3AI-type intermetallic ribbons

An aluminum matrix composite containing rapidly solidified Ni₇₅Al₂₃B₁Zr₁ (at. pct) ribbons has been fabricated by casting at 700 °C, 715 °C, 730 °C, and 875 °C. Microstructural investigation has shown that the matrix contains particles with a composition between Al₃Ni and eutectic. The interfacial zones composed of several layers with different aluminum and nickel contents are observed around the ribbons. The sequence of layers from the ribbon outward in the specimens fabricated at 700 °C, 715 °C, and 730 °C is as follows: AINi → Al₃Ni₂ → the outer layer between Al₃Ni and eutectic. Composite specimens fabricated at 875 °C contain two types of interfacial zones: a single-layer AINi and a triple-layer zone. The first two layers in the triplelayer zone are exactly the same as their counterparts in the specimens fabricated at lower temperatures. The outer layer has a composition close to the Al₃Ni compound. The thickness of the AINi layer increases continuously with the increasing casting temperature. Within the experimental error, the thickness of the Al₃Ni₂ layer seems to be independent of casting temperature. The thickness of the outer layer in the specimens fabricated at 700 °C to 730 °C (Al₃Ni plus eutectic) increases with the casting temperature. However, the outer layer in the 875 °C specimen (Al₃Ni) is much thinner than the others.     

Oxidation Processes for Alloys in the Ni – Zr System. Part 2. Phase Composition of Scale Formed on Ni7Zr2

We have used x-ray and metallographic layer-by-layer phase analysis to study the structure and composition of scale formed on the alloy Ni₇Zr₂ during its oxidation in air over a period of 1 h and 10 h in the temperature range 500-1200°C. In the scale we find NiO, the cubic and monoclinic modifications of ZrO₂, and also Ni and Ni₅Zr. The phase components are nonuniformly distributed over the thickness of the scale. The outer scale consists of the oxides NiO and ZrO₂, while the composition of the inner scale includes Ni and Ni₅Zr in addition to monoclinic ZrO₂. Cubic ZrO₂ is formed on the surface of the specimen in the initial stages of its oxidation at 500-700°C. For T ≥ 900°C, on the surface of the scale we find both modifications of ZrO₂, while the nickel phase is itself a solid solution Ni(Zr). We note that the mechanisms for the formation of low-temperature (T ≤ 800°C) and high-temperature (T ≥ 900°C) scales are different. It is hypothesized that these differences are determined mainly by the fact that at high temperatures, diffusion of zirconium ions toward the outer boundary of the scale is superimposed on diffusion of oxygen toward the scale – alloy boundary.     

Study of Aging-Induced Degradation of Fracture Resistance of Alloy 617 Toward High-Temperature Applications

For the Alloy 617, the effect of aging on the fracture energy degradation has been investigated after aging for different time periods at 1023 K (750 °C). A sharp reduction in impact energy (by ~55 pct vis-à-vis the as-received material) after 1000 hours of aging, as evaluated from room-temperature Charpy impact tests, has been observed. Further aging up to 10,000 hours has led to a degradation of fracture energy up to ~78 pct. Fractographic examinations using scanning electron microscopy (SEM) have revealed a change in fracture mode from fibrous-ductile for the un-aged material to intergranular mode for the aged one. The extent of intergranular fracture increases with the increasing aging time, indicating a tendency of the material to undergo grain boundary embrittlement over long-term aging. Analysis of the transmission electron microscopy (TEM) micrographs along with selected area diffraction (SAD) patterns for the samples aged at 10,000 hours revealed finely dispersed γ′ precipitates of size 30 to 40 nm, rich in Al and Ti, along with extensive precipitation of M₂₃C₆ at the grain boundaries. In addition, the presence of Ni₃Si of size in the range of 110 to 120 nm also has been noticed. The extensive precipitation of M₂₃C₆ at the grain boundaries have been considered as a major reason for aging-induced embrittlement of this material.     

The effect of ordering on the hydrogen embrittlement susceptibility of Ni2Cr

Isothermal annealing at 500 °C for various lengths of time after rapid quenching from 900 °C results in different degrees of ordering in Ni₂Cr. Tensile specimens of disordered and ordered Ni₂Cr were subjected to hydrogen embrittling and nonembrittling environments before and during tensile testing. The hydrogen embrittlement susceptibility was determined by reduction-in-area losses (RA loss) after failure. The results of the tensile tests indicated that the disordered alloy and highly ordered alloy were the most susceptible to hydrogen embrittlement. The RA loss in each case was approximately 70 pct. The tests revealed a minimum susceptibility to hydrogen embrittlement between 40 and 50 pct order. The fracture surfaces of the specimens, as examined by scanning electron microscopy, showed that the extent of embrittlement is correlated with the amount of intergranular failure. The mechanisms for failure in Ni₂Cr appear to depend upon the extent of aging in the alloy.     

Structure and mechanical properties of boron-doped cubic zirconium trialuminides

Cubic (L1₂) ternary zirconium trialuminides macroalloyed with Cu(Al₅CuZr₂), Mn(Al₆₆Mn₉Zr₂₅), and Cr(Al₆₇Cr₈Zr₂₅) (atomic percent) and doped with 50 and 100 ppm boron were fabricated by induction melting. Their as-cast microstructures are characterized by a small amount of porosity (1 to 2 pct) and second phase (2 to 3 pct). Boron seems to slightly enhance porosity (up to 3.3 pct) in Al₅CuZr₂ +100 ppm B alloy, and it also promotes some compositional inhomo-geneity in Al₆₆Mn₉Zr₂₅ alloy. Vickers microhardness and compressive properties at room temperature (RT), peak strength temperature (500 °C to 600 °C) and 900 °C were investigated. Microcracking development was also investigated in Al₅CuZr₂ +100 ppm boron alloy exhibiting a stepped load-deflection curve. Vickers microhardness strongly depends on load, similarly to boron-free cubic ternary zirconium and titanium trialuminides, and increases in a systematic way with increasing boron content which seems to indicate a solid solution strengthening effect. At RT, 0.2 pct offset yield strength is not increased by the boron doping in most of the alloys studied except for Al₆₆Mn₉Zr₂₅ + 50 ppm B alloy. Permanent deformation (apparent ductility) at ultimate compressive strength is not enhanced by boron doping. In Al₅CuZr₂ +100 ppm B alloy microcracks start nucleating and proliferating in the elastic region of load-deflection curve in characteristic “bursts” accompanied by a “click” sound and the appearance of a discernible step on the load-deflection curve. Pre-existing pores are observed to be active centers of microcracking.     

Mechanical alloying of IN-738

Mechanical alloying has been applied to produce a dispersion-strengthened superalloy IN-738 containing 1.5 wt pct Y₂O₃. Annealing of extrusion bars above the recrystallization temperature of 1160°C can be described by three stages of recrystallization:finegrain; isotropic coarse-grain; and fibrous coarse grain growth. A maximum grain length of 550 μm and a maximum grain aspect ratio of 4.8 have been obtained for an alloy, which had been extruded at 1100°C and annealed at 1280°C and 1270°C for 3 h, respectively. The three stages of grain growth are explained in terms of recovery, differences in nucleation rate and dispersoid concentration in the two normal directions and release in stored cold work. Secondary recrystallization can be excluded as a mechanism for fibrous grain coarsening. Dispersion-strengthened IN-738, heat treated to a coarse elongated grain structure, has both high intermediate temperature strength and high elevated temperature strength. The creep strength at 1000°C exceeds that of cast or directionally solidified IN-738 after 300 h service life. The failure mechanism at elevated temperature is intergranular fracture along transverse grain boundaries, nucleated by cavities that form during grain boundary sliding. Nucleation of voids is retarded in the creep specimens due to diffusional accommodation of grain boundary sliding. A depletion of surface zones of chromium, aluminum and titanium contributes to initiation of creep failure at 1000°C.     

Mechanical alloying of IN-738

Mechanical alloying has been applied to produce a dispersion-strengthened superalloy IN-738 containing 1.5 wt pct Y₂O₃. Annealing of extrusion bars above the recrystallization temperature of 1160°C can be described by three stages of recrystallization:finegrain; isotropic coarse-grain; and fibrous coarse grain growth. A maximum grain length of 550 μm and a maximum grain aspect ratio of 4.8 have been obtained for an alloy, which had been extruded at 1100°C and annealed at 1280°C and 1270°C for 3 h, respectively. The three stages of grain growth are explained in terms of recovery, differences in nucleation rate and dispersoid concentration in the two normal directions and release in stored cold work. Secondary recrystallization can be excluded as a mechanism for fibrous grain coarsening. Dispersion-strengthened IN-738, heat treated to a coarse elongated grain structure, has both high intermediate temperature strength and high elevated temperature strength. The creep strength at 1000°C exceeds that of cast or directionally solidified IN-738 after 300 h service life. The failure mechanism at elevated temperature is intergranular fracture along transverse grain boundaries, nucleated by cavities that form during grain boundary sliding. Nucleation of voids is retarded in the creep specimens due to diffusional accommodation of grain boundary sliding. A depletion of surface zones of chromium, aluminum and titanium contributes to initiation of creep failure at 1000°C.     

Heat treatment of 706 alloy for optimum 1200 °F stress-rupture properties

Evaluation of a commercial heat treatment for 706 alloy indicated that it resulted in relatively low 1200° F stress rupture ductility. It was determined that this was caused by a solution treatment which dissolved all of the age-hardening phases in the alloy and caused a coarse grain size and supersaturated matrix condition. Based upon extensive fine structure study of the 706 alloy as well as previous experience with 718 alloy and other Fe?Ni-base superalloys, a heat treatment is developed which effectively optimizes the 1200°F stress-rupture properties of the alloy. The key to best properties was found to be the precipitation of globular to plate-like Ni₃Cb/Ni₃Ti at the grain boundaries in conjunction with maintaining a fine as-forged grain structure.     

Localized corrosion susceptibility of Al-Li-Cu-Mg-Zn alloy AF/C458 due to interrupted quenching from solutionizing temperature

Isothermal time-temperature-localized corrosion-behavior curves were determined for the Al-1.8Li-2.70Cu-0.6Mg-0.3Zn alloy AF/C458, to understand the effect of slow or delayed quenching on localized corrosion susceptibility. Alloy samples were subject to a series of systematic interrupted quenching experiments conducted at temperatures ranging from 480 °C to 230 °C for times ranging from 5 to 1000 seconds. Individual samples were then exposed to an oxidizing aqueous chloride solution consisting of 57 g/L NaCl plus 10 mL/L H₂O₂ to induce localized attack. The localized corrosion mode was characterized by optical microscopy. Additionally, the microstructure of selected samples was characterized by transmission electron microscopy (TEM) to relate the corrosion mode and morphology to microstructural features. Results showed that only pitting attack was exhibited by samples subjected to isothermal treatment at temperatures greater than 430 °C. At temperatures ranging from 280 °C to 430 °C, isothermal treatment tended to induce susceptibility to intergranular attack (IGA) and intersubgranular attack (ISGA) for all treatment times investigated. For isothermal treatments at temperatures lower than 280 °C, only pitting was observed for treatment times less than about 30 seconds, while IGA and ISGA were observed for longer treatment times. Comparisons showed that the time-temperature domains for IGA and ISGA were virtually coincident. Based on this finding and the results from TEM characterization, IGA and ISGA appear to be related to the precipitation of a Zn-modified T ₁ (Al₂(Cu,Zn)Li) precipitate, which can occur both on low-angle and high-angle grain boundaries in this alloy. When the alloy is resistant to IGA and ISGA, the grain boundaries are decorated by θ′ (Al₂Cu), and T _B (Al₇Cu₄Li) phase particles, or subgrain boundaries are populated by a comparatively low density of T ₁ precipitates. It is, therefore, speculated that θ′ and T _B are more corrosion-resistant precipitate phases than T ₁, and that a critical concentration of boundary T ₁ must exist for IGA or ISGA to occur.     

Evolution of Strain and Microstructure During Creep of Wrought AE42 and ZE10 Magnesium Alloys

Wrought magnesium alloys have been extensively used in the aerospace, electronics and automotive industries, where component weight is of concern and ambient temperatures remain below 100 °C. Undesirable creep relaxation of the wrought alloys above this temperature has been generally attributed to grain boundary sliding and plastic deformation leading to intergranular failure. The objective of this study was to investigate the compressive creep performance and microstructure of two wrought magnesium alloys (AE42 and ZE10) developed for high temperature applications. The total deformation of the AE42 and ZE10 alloys was 2.4 and 0.2 %, respectively, after 24 h creep test at 175 °C and 50 MPa. The poor creep performance of the AE42 alloy was explained via neutron diffraction studies which revealed that the elastic compressive response of $ (10\bar{1}0),\;(10\bar{1}1)\;{\text{and}}\;(2\bar{1}\bar{1}0) $ planes was significantly more anisotropic in the AE42 than in the ZE10 alloy. Further, microstructural analysis revealed ~10 % increase in grain size due to creep, with additional $ (10\bar{1}2) $ and $ (11\bar{2}1) $ twinning in the AE42 alloy. Precipitation of β-Mg₁₇Al₁₂ phase in the AE42 alloy possibly contributed to grain boundary sliding and high plastic strain during creep testing.     

The influence of grain boundary precipitation on the measurement of chromium redistribution and phosphorus segregation in Ni-16Cr-9Fe

The redistribution of chromium at the grain boundary and the segregation of phosphorus to the grain boundary in Ni-16Cr-9Fe is measured following thermal treatment at 700 °C for 1 to 100 hours. The addition of carbon to the base alloy results in the formation of Cr₇C₃ precipitates at the grain boundary and the formation of a chromium depleted zone in the adjacent matrix. Measurement of the Cr concentration is affected by the presence of Cr-rich carbides, and a technique of ratioing the Auger signal of the element of interest to a sum of the signals of elements present in the carbide and the matrix is required to minimize the scatter in the data. The presence of carbides does not affect the kinetics or extent of phosphorus segregation to the grain boundary, and there is no evidence of co-segregation of phosphorus with any major alloying element. The free energy of segregation of phosphorus is determined to be 46.2 KJ/mole at 1100 °C and 40.8 KJ/mole at 700 °C. Results show that the intergranular fracture path is along the carbide-matrix interface as opposed to through the carbides or some distance into the matrix. These results permit the calculation of the coverage of the grain boundary with carbides.