Microstructure and phase identification in type 304 stainless steel-zirconium alloys

Stainless steel-zirconium alloys have been developed at Argonne National Laboratory to contain radioactive metal isotopes isolated from spent nuclear fuel. This article discusses the various phases that are formed in as-cast alloys of type 304 stainless steel and zirconium that contain up to 92 wt pct Zr. Microstructural characterization was performed by scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDS), and crystal structure information was obtained by X-ray diffraction. Type 304SS-Zr alloys with 5 and 10 wt pct Zr have a three-phase microstructure—austenite, ferrite, and the Laves intermetallic, Zr(Fe,Cr,Ni)_2+x. whereas alloys with 15, 20, and 30 wt pct Zr contain only two phases—ferrite and Zr(Fe,Cr,Ni)_2+x. Alloys with 45 to 67 wt pct Zr contain a mixture of Zr(Fe,Cr,Ni)_2+x and Zr₂(Ni,Fe), whereas alloys with 83 and 92 wt pct Zr contain three phases—α-Zr, Zr₂(Ni,Fe), and Zr(Fe,Cr,Ni)_2+x. Fe₃Zr-type and Zr₃Fe-type phases were not observed in the type 304SS-Zr alloys. The changes in alloy microstructure with zirconium content have been correlated to the Fe-Zr binary phase diagram.     

Microstructure and mechanical properties of ternary intermetallic compound dispersed P/M Al-Mn-X-Zr (x=Cu,Ni) alloys

Heat-resistant aluminum alloys are generally developed by dispersing stable intermetallic compounds by adding transition metals (TM) whose diffusion coefficient in aluminum alloys is low even at high temperatures. Commonly used intermetallic compounds include Al-TM binary intermetallic compounds, for example, Al₆Fe, Al₃Ti and Al₃Ni. By contrast, multicomponent intermetallic compounds are hardly used. The present study focuses on Al-Mn-Cu and Al-Mn-Ni ternary intermetallic compounds, and by finely dispersing these intermetallic compounds, attempts to develop heat-resistant alloys. Through the atomization method, Al-(4.96–5.96)Mn-(6.82–7.53)Cu-0.4Zr and Al-(5.48–8.76)Mn-(2.23–4.32)Ni-0.4Zr (in mass%) powders were fabricated, and by degassing these powders at 773 K, intermetallic compounds were precipitated. These powders were then solidified into extrudates by hot extrusion at 773 K. The microstructural characterization of powders and exrudates was carried out by XRD analysis, SEM/EDX and TEM. The mechanical properties of extrudates were determined at room temperature, 523 K and 573 K. In Al-Mn-Cu alloys, while a small amount of Al₂Cu was crystallized, precipitated Al₂₀Mn₃Cu₂ intermetallic compounds were mainly dispersed. In Al-Mn-Ni alloys, while a small amount of Al₆Mn intermetallic compounds was precipitated, the precipitated A₆₀Mn₁₁Ni₄ intermetallic compounds were mainly dispersed. Both ternary intermetallic compounds were about 200 nm in size. The compounds were elliptical, and their longitudinal direction was oriented along the extrusion direction. In the Al-Mn-Cu alloys, since the work hardening at room temperature was high, the tensile strength became 569 MPa. At elevated temperatures, since hardly any work hardening was observed, the tensile strength decreased markedly. However, in Al-Mn-Ni alloys, since the work hardening is low even at room temperature, the roomtemperature strength is not high. Thus, the decrease in tensile strength at elevated temperatures is relatively small and a high strength was obtained at 523 K and 573 K: 276 MPa and 207 MPa, respectively.     

Microstructural Evolution of the Interface Between Pure Titanium and Low Melting Point Zr-Ti-Ni(Cu) Filler Metals

Low melting point Zr-based filler metals with melting point depressants (MPDs) such as Cu and Ni elements are used for titanium brazing. However, the phase transition of the filler metals in the titanium joint needs to be explained, since the main element of Zr in the filler metals differs from that of the parent titanium alloys. In addition, since the MPDs easily form brittle intermetallics, that deteriorate joint properties, the phase evolution they cause needs to be studied. Zr-based filler metals having Cu content from 0 to 12 at. pct and Ni content from 12 to 24 at. pct with a melting temperature range of 1062 K to 1082 K (789 °C to 809 °C) were wetting-tested on a titanium plate to investigate the phase transformation and evolution at the interface between the titanium plate and the filler metals. In the interface, the alloys system with Zr, Zr₂Ni, and (Ti,Zr)₂Ni phases was easily changed to a Ti-based alloy system with Ti, Ti₂Ni, and (Ti,Zr)₂Ni phases, by the local melting of parent titanium. The dissolution depths of the parent metal were increased with increasing Ni content in the filler metals because Ni has a faster diffusion rate than Cu. Instead, slow diffusion of Cu into titanium substrate leads to the accumulation of Cu at the molten zone of the interface, which could form undesirable Ti _x Cu _y intermetallics. This study confirmed that Zr-based filler metals are compatible with the parent titanium metal with the minimum content of MPDs.     

Effect of Bonding Time on Interfacial Reaction and Mechanical Properties of Diffusion-Bonded Joint Between Ti-6Al-4V and 304 Stainless Steel Using Nickel as an Intermediate Material

In the current study, solid-state diffusion bonding between Ti-6Al-4V (TiA) and 304 stainless steel (SS) using pure nickel (Ni) of 200-μm thickness as an intermediate material was carried out in vacuum. Uniaxial compressive pressure and temperature were kept at 4 MPa and 1023 K (750 °C), respectively, and the bonding time was varied from 30 to 120 minutes in steps of 15 minutes. Scanning electron microscopy images, in backscattered electron mode, revealed the layerwise Ti-Ni-based intermetallics like either Ni₃Ti or both Ni₃Ti and NiTi at titanium alloy-nickel (TiA/Ni) interface, whereas nickel-stainless steel (Ni/SS) interface was free from intermetallic phases for all the joints. Chemical composition of the reaction layers was determined by energy dispersive spectroscopy (SEM–EDS) and confirmed by X-ray diffraction study. Maximum tensile strength of ~382 MPa along with ~3.7 pct ductility was observed for the joints processed for 60 minutes. It was found that the extent of diffusion zone at Ni/SS interface was greater than that of TiA/Ni interface. From the microhardness profile, fractured surfaces, and fracture path, it was demonstrated that the failure of the joints was initiated and propagated apparently at TiA/Ni interface near Ni₃Ti intermetallic for bonding time less than 90 minutes, and through Ni for bonding time 90 minutes and greater.     

Synthesis of Invar 36 type alloys from elemental and prealloyed powders by mechanical alloying

Invar 36 (Fe₆₄Ni₃₆) nanocrystalline powders were successfully obtained by the mechanical alloying process. The mechanically alloyed Invar 36 powders were obtained from both, Fe–Ni elemental and Fe–Ni₃Fe prealloyed powders. XRD, DSC and magnetic measurements were used to characterise the Invar 36 powders. The lattice parameter evolution versus temperature of Invar 36 powders was investigated by in-situ high-temperature X-ray diffraction (HT-XRD). For both, Invar 36 (Fe, Ni) and Invar 36 (Fe, Ni₃Fe) powders, the lattice parameter values are constant up to about 350°C. The magnetic measurement also indicated that the Invar 36-type alloys are formed after 16?h of milling.     

Effect of Bonding Temperature on Interfacial Reaction and Mechanical Properties of Diffusion-Bonded Joint Between Ti-6Al-4V and 304 Stainless Steel Using Nickel as an Intermediate Material

An investigation was carried out on the solid-state diffusion bonding between Ti-6Al-4V (TiA) and 304 stainless steel (SS) using pure nickel (Ni) of 200-μm thickness as an intermediate material prepared in vacuum in the temperature range from 973 K to 1073 K (700 °C to 800 °C) in steps of 298 K (25 °C) using uniaxial compressive pressure of 3 MPa and 60 minutes as bonding time. Scanning electron microscopy images, in backscattered electron mode, had revealed existence of layerwise Ti-Ni-based intermetallics such as either Ni₃Ti or both Ni₃Ti and NiTi at titanium alloy-nickel (TiA/Ni) interface, whereas nickel-stainless steel (Ni/SS) diffusion zone was free from intermetallic phases for all joints processed. Chemical composition of the reaction layers was determined in atomic percentage by energy dispersive spectroscopy and confirmed by X-ray diffraction study. Room-temperature properties of the bonded joints were characterized using microhardness evaluation and tensile testing. The maximum hardness value of ~800 HV was observed at TiA/Ni interface for the bond processed at 1073 K (800 °C). The hardness value at Ni/SS interface for all the bonds was found to be ~330 HV. Maximum tensile strength of ~206 MPa along with ~2.9 pct ductility was obtained for the joint processed at 1023 K (750 °C). It was observed from the activation study that the diffusion rate at TiA/Ni interface is lesser than that at the Ni/SS interface. From microhardness profile, fractured surfaces and fracture path, it was demonstrated that failure of the joints was initiated and propagated apparently at the TiA/Ni interface near Ni₃Ti intermetallic phase.     

The microstructure of an Fe-Mn-Si-Cr-Ni stainless steel shape memory alloy

The microstructure and phase stability of the Fe-15Mn-7Si-9Cr-5Ni stainless steel shape memory alloy in the temperature range of 600 °C to 1200 °C was investigated using optical and transmission electron microscopy, X-ray diffractometry (XRD), differential scanning calorimetry (DSC), and chemical analysis techniques. The microstructural studies show that an austenite single-phase field exists in the temperature range of 1000 °C to 1100 °C, above 1100 °C, there exists a three-phase field consisting of austenite, δ-ferrite, and the (Fe,Mn)₃Si intermetallic phase; within the temperature range of 700 °C to 1000 °C, a two-phase field consisting of austenite and the Fe₅Ni₃Si₂ type intermetallic phase exists; and below 700 °C, there exists a single austenite phase field. Apart from these equilibrium phases, the austenite grains show the presence of athermal ɛ martensite. The athermal α′ martensite has also been observed for the first time in these stainless steel shape memory alloys and is produced through the γ-ɛ-α′ transformation sequence.     

Examination of Multiphase (Zr,Ti)(V,Cr,Mn,Ni)<Subscript>2</Subscript> Ni-MH Electrode Alloys: Part II. Solid-State Transformation of the Interdendritic B2 Phase

Solidification microstructure of multicomponent (Zr,Ti)-Ni-(V,Cr,Mn,Co) alloys intended for use as negative electrodes in Ni-metal hydride (Ni-MH) batteries was studied in Part I of this series of articles. Part II of the series examines the complex internal structure of the interdendritic grains formed by solid-state transformation and believed to play an important role in the electrochemical charge/discharge characteristics of the overall alloy composition. By studying one alloy, Zr₂₁Ti_12.5V₁₀Cr_5.5Mn_5.1Co_5.0Ni_40.2Al_0.5Sn_0.3, it is shown that the interdendritic grains solidify as a B2 (Ti,Zr)₄₄(Ni,TM)₅₆ phase, and then undergo transformation to Zr₇Ni₁₀-type, Zr₉Ni₁₁-type, and martensitic phases. The transformations obey orientation relationships between the high-temperature B2 phase and the low-temperature Zr-Ni–type intermetallics, and consequently lead to a multivariant structure. The major orientation relationship for the orthorhombic Zr₇Ni₁₀ type is [011]_Zr7Ni10//[001]_B2; (100)_Zr7Ni10//(100)_B2. The orientation relationship for the tetragonal Zr₉Ni₁₁ type is [001]_Zr9Ni11//[001]_B2; (130)_Zr9Ni11//(100)_B2. Binary Ni-Zr and ternary Ti-Ni-Zr phase diagrams were used to rationalize the formation of the observed domain structure.     

Overview of Strategies for High-Temperature Creep and Oxidation Resistance of Alumina-Forming Austenitic Stainless Steels

A family of creep-resistant, alumina-forming austenitic (AFA) stainless steel alloys is under development for structural use in fossil energy conversion and combustion system applications. The AFA alloys developed to date exhibit comparable creep-rupture lives to state-of-the-art advanced austenitic alloys, and superior oxidation resistance in the ~923 K to 1173 K (650 °C to 900 °C) temperature range due to the formation of a protective Al₂O₃ scale rather than the Cr₂O₃ scales that form on conventional stainless steel alloys. This article overviews the alloy design approaches used to obtain high-temperature creep strength in AFA alloys via considerations of phase equilibrium from thermodynamic calculations as well as microstructure characterization. Strengthening precipitates under evaluation include MC-type carbides or intermetallic phases such as NiAl-B2, Fe₂(Mo,Nb)-Laves, Ni₃Al-L1₂, etc. in the austenitic single-phase matrix. Creep, tensile, and oxidation properties of the AFA alloys are discussed relative to compositional and microstructural factors.     

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 Effect of Aluminum Additions on the Microstructure and Thermomechanical Behavior of NiTiZr Shape-Memory Alloys

The microstructure, thermal cycling, and mechanical behavior of Ni_48.5Ti_31.5?x Zr₂₀Al _x (x?=?0, 1, 2, 3) alloys were studied in the solution-treated and aged condition using microscopy techniques, differential scanning calorimetry, and compression tests. The microscopy techniques used include optical, scanning, and transmission electron microscopy, and three-dimensional, atom?Cprobe microscopy. The results of this study indicated a strong dependence of the transformation behavior on alloy chemistry and thermal cycling. The aluminum additions served to decrease transformation behaviors from 351?K to 596?K (78?°C to 323?°C) and reduce thermal stability. Additionally, aluminum was shown to increase the plateau stress in the aged condition, whereas the formation of coarse-grained intermetallic phases caused the embrittlement of the microstructure, reducing its ductility. The addition of Al resulted in the refinement of the coarse, lenticular precipitates identified as Ni₄(Ti,Zr)₃.     

Understanding the high-temperature deformation

Engineering, University of Texas at Austin, Austin, TX 78712 While much of the high-temperature intermetallics research has centered around Ni3Al and other aluminum-based systems, the present study focuses on the Engel-Brewer Ll₂ intermetallic Ir₃Zr, which has a melting temperature approaching that of ceramics (2280 °C). Due to limited material availability, the technique of microindentation was used to study both the temperature and time dependence of strength. Because of the widely held belief that certain mechanical properties of intermetallics scale roughly with temperature, Ir₃Zr was expected to exhibit high strength. The microhardness was observed to vary from 225 MPa at room temperature to 75 MPa at 1400 °C, which is significantly lower than the behavior of Ni₃Al. The activation energy for creep was determined to be 467 kJ/mole, and the stress exponent was found to be 18.2. The ordering energy of this system was calculated to be 0.114 eV. If it can be assumed that high ordering energy correlates to a high antiphase boundary (APB) energy, then the behavior of this system is consistent with a model that predicts highly glissile dislocation cores.     

Effect of alloying elements on martensitic transformation in the binary NiAl(β) phase alloys

The characteristics of the B2(β) to L1₀(β′) martensitic transformation in NiAl base alloys containing a small amount of third elements have been investigated by differential scanning calorimetry (DSC), X-ray diffraction (XRD), and transmission electron microscopy (TEM). It is found that in addition to the normal Ll₀ (3R) martensite, the 7R martensite is also present in the ternary alloys containing Ti, Mo, Ag, Ta, or Zr. While the addition of third elements X (X: Ti, V, Cr, Mn, Fe, Zr, Nb, Mo, Ta, W, and Si) to the binary Ni₆₄Al₃₆ alloy stabilizes the parentβ phase, thereby lowering the M_s temperature, addition of third elements such as Co, Cu, or Ag destabilizes theβ phase, increasing the M_s temperature. The occurrence of the 7R martensite structure is attributed to solid solution hardening arising from the difference in atomic size between Ni and Al and the third elements added. The variation in M_s temperature with third element additions is primarily ascribed to the difference in lattice stabilities of the bcc and fcc phases of the alloying elements.     

Intermetallic diffusion coatings for enhanced hot-salt oxidation resistance of nitrogen-containing austenitic stainless steels

This article presents the preparation, characterization, and hot-salt oxidation behavior of nitrogen-containing type 316L stainless steel (SS), surface modified with intermetallic coatings. Three different types of intermetallic coating systems, containing aluminum, titanium, and titanium/aluminum multilayers, were formed by diffusion annealing of type 316L austenitic SS containing 0.015, 0.1, 0.2, and 0.56 pct nitrogen. Analysis by using X-ray diffraction (XRD), scanning electron microscopy (SEM), and secondary ion mass spectroscopy (SIMS) confirmed the formation of various intermetallic phases such as AIN, Al₁₃Fe₄, FeAl₂, FeTi, Ti₂N, and Ti₃Al in the coatings. Hot salt oxidation behavior of the uncoated and surface-modified stainless steels was assessed by periodic monitoring of the weight changes of NaCl salt-applied alloys kept in an air furnace at 1023 K up to 250 hours. The oxide scales formed were examined by XRD and stereomicroscopy. Among the various surface modifications investigated in the present study, the results indicate that the titanium-modified alloys show the best hot-salt oxidation resistance with the formation of an adherent, protective, thin, and continuous oxide layer. Among the four N-containing alloys investigated, the titanium and Ti/Al multilayer modified 0.56 pct N alloy showed the best hot-salt oxidation resistance as compared to uncoated alloys. The slower corrosion kinetics and adherent scale morphology indicate that the surface-modified titanium intermetallic coatings could provide high-temperature service applications up to 1073 K, particularly in chloride containing atmospheres, for austenitic stainless steels.     

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).     

Structural and Magnetic Properties of Single Domain PEG-Coated Fe<Subscript>x</Subscript>Ni<Subscript>80−x</Subscript>Co<Subscript>20</Subscript> (x = 20, 40) Ternary Alloys Synthesized by Chemical Method

In the present study, we have synthesized nanocrystalline Fe _x Ni_80?xCo₂₀ (x = 20 and 40) alloys via modified sodium borohydride reduction route in aqueous medium. The phase purity and microstructural analysis of the materials were done using powder X-ray diffraction (XRD) and transmission electron microscopy (TEM). Fe–Ni–Co alloys crystallize in face centered cubic (fcc) structure having the values of lattice parameters, a = 3.5460 and 3.5634 Å, for Fe₂₀Ni₆₀Co₂₀ and Fe₄₀Ni₄₀Co₂₀, respectively. The average crystallite sizes were found to be 30 and 26 nm for Fe₂₀Ni₆₀Co₂₀ and Fe₄₀Ni₄₀Co₂₀, respectively. Fe₂₀Ni₆₀Co₂₀ and Fe₄₀Ni₄₀Co₂₀ alloys crystallize in nearly spherical shape morphologies having TEM particle sizes of 24 and 36 nm, respectively. Selected area electron diffraction (SAED) patterns confirm face centered cubic crystal structure and polycrystalline nature of the materials. Low temperature studies indicate that Fe–Ni–Co alloys show ferromagnetism at least up to 300 K. The effective magnetic anisotropy constant (K) for Fe₂₀Ni₆₀Co₂₀ and Fe₄₀Ni₄₀Co₂₀ alloys are found to be 14 and 4.2 kJm^?3, respectively. These results have been explained on the basis of size, shape and surface effects in the nano-materials. The contribution corresponding to spin relaxation and dipolar interactions between nano-particles on magnetic characteristics of Fe₂₀Ni₆₀Co₂₀ and Fe₄₀Ni₄₀Co₂₀ alloys have been explained qualitatively during the magnetic analysis.     

Overview on Basic Types of Hot Rolling Textures of Steels

The present study gives a review on basic types of crystallographic textures developing during hot rolling of polycrystalline steels. The results are grouped into three fundamental classes of textures. The first group comprises pure Fe, some weakly bonded B2 and DO₃ structured intermetallics, as well as closely related alloys such as ferritic low carbon and microalloyed interstitial free steels. The second group includes highly alloyed corrosion‐resistant ferritic stainless and Fe‐Si transformer steels. Typical examples are steels with about 10 wt.%‐17 wt.% Cr, with about 3 wt.% Si, as well as body centered cubic transition metals such as Ta, Mo, and Nb which do not undergo any phase transformation during hot rolling. The third group comprises stable and instable austenitic stainless steels for instance on the basis of larger amounts of Cr and Ni or on Mn as well as duplex steels. Most L1₂ structured intermetallic alloys can also be assigned to this group. The suggested classification scheme is discussed in terms of different processing parameters, thermodynamics, microstructure, and crystallographic aspects.     

Structural aspects of the manufacture of semiproducts made from titanium nickelide-based alloys

The effects of the charge composition, casting method, and metal forming method on the structure and shape-memory-effect (SME) and superelasticity characteristics of titanium nickelide-based alloys are studied. The shape recovery temperatures of semiproducts are shown to depend substantially on the volume fraction of the Ti₂Ni intermetallic phase, whose formation is stimulated by the oxygen present in a charge or absorbed during casting. An increase in the volume fraction of Ti₂Ni in an alloy leads to nickel enrichment of the B2 phase and a decrease in the shape recovery temperatures. Subsequent metal forming at the stage of semiproduct manufacture only weakly affects the volume fraction of Ti₂Ni and favors the formation of its equiaxed shape and a more uniform Ti₂Ni distribution in the B2 matrix. In alloys where the B2 phase contains more than 56.5 wt % Ni, quenching from temperatures above 600°C and aging in the temperature range 400–500°C result in the dissolution and precipitation, respectively, of the nickel-rich Ti₃Ni₄ and Ti₂Ni₃ intermetallics. Therefore, the shape recovery temperatures of semiproducts and finished products can be controlled. Moreover, as the aging temperature changes, the volume fraction and size of nickel-rich intermetallic particles, the slip stresses, and the SME force characteristics change. For example, to increase the compression forces for osteosynthesis fixation devices, one has to use a titanium nickelide-based alloy with a high nickel content in the B2 phase and to perform aging at low temperatures (400–450°C).     

Investigations on hydrogen storage properties of LaMg8.52Ni2.23M0.15 （M=Ni, Cu, Cr） alloys

LaMg8.52Ni2.23M0.15 (M=Ni, Cu, Cr) alloys were prepared by induction melting. X-ray diffraction showed that all the three alloys had a multiphase structure, consisting of La2Mg17, LaMg2Ni and Mg2Ni phases. Energy dispersive X-ray spectrometer results revealed that most of Cu and Cr distributed in Mg2Ni phase. La2Mg17 and LaMg2Ni phases decomposed into MgH2, Mg2NiH4 and LaH3 phases during the hydrogenation process. Hydriding/dehydriding measurements indicated that the reversible hydrogen storage capacities of Mg2Ni phase in LaMg8.52Ni2.23M0.15 (M=Cu, Cr) alloys increased to 1.05 wt.% and 0.97 wt.% from 0.79 wt.% of Mg2Ni phase in LaMg8.52Ni2.38 alloy at 523 K. Partial substitution of Cu and Cr for Ni decreased the onset dehydrogenation temperature of the alloy hydrides and the temperature lowered by 18.20 and 5.50 K, respectively. The improvement in the dehydrogenation property of the alloys was attributed to that Cu and Cr decreased the stability of Mg2NiH4 phase.