Rare Earth-rich Phase Research Articles

Structure and magnetostriction of Sm 1− xPr xFe 2 and Sm 0.9Pr 0.1(Fe 1− yB y) 2 alloys

The structure, magnetostriction and anisotropy of Sm 1− x Pr x Fe 2 and Sm 0.9Pr 0.1(Fe 1− y B y ) 2 alloys have been investigated using X-ray diffraction, standard strain gauge technique and Mössbauer spectrum. It is found that the matrix of homogenized Sm 1− x Pr x Fe 2 alloys is the (Sm,Pr)Fe 2 phase with MgCu 2-type cubic structure and the minor phases are (Sm,Pr)Fe 3 phase and rare earth-rich phase when x⩽0.3. When x=0.4, the (Sm,Pr)Fe 3 phase is found to be the main phase and the minor phases become (Sm,Pr)Fe 2, (Sm,Pr) 2Fe 17 and rare earth-rich phase. With further increasing Pr content, the amount of minor phases, (Sm,Pr) 2Fe 17 and rare earth-rich phase increase, but that of (Sm,Pr)Fe 2 remains almost unchanged. For Sm 0.9Pr 0.1(Fe 1− y B y ) 2 alloys, the (Sm,Pr)(Fe,B) 2 as a matrix is up to y=0.2, and the small amount of second phase is a rare earth-rich phase when y>0.05. The magnetostriction of Sm 1− x Pr x Fe 2 alloys exhibits a peak near x=0.1 and that of Sm 0.9Pr 0.1(Fe 1− y B y ) 2 alloys decreases with increasing B content. The Mössbauer spectrum analysis shows that the easy axis of magnetization for Sm 1− x Pr x Fe 2 alloys lies in the 〈1 1 1〉 direction up to x=0.3 and the anisotropy of Sm 1− x Pr x Fe 2 alloys decreases with increasing Pr content.

Structure and magnetic and magnetostrictive properties of (Tb/sub 0.7/Dy/sub 0.3/)/sub 0.7/Pr/sub 0.3/(Fe/sub 1-x/Co/sub x/)/sub 1.85/ (0≤×≤0.6)

Measurements of the structure, magnetic properties, and magnetostriction were made on arc-melted polycrystalline (Tb/sub 0.7/Dy/sub 0.3/)/sub 0.7/Pr/sub 0.3/(Fe/sub 1-x/Co/sub x/)/sub 1.85/ (0/spl les//spl times//spl les/0.6) alloys by X-ray diffraction, initial susceptibility, vibrating sample magnetometry, and standard strain gauge techniques. The measurements showed that the matrix of (Tb/sub 0.7/Dy/sub 0.3/)/sub 0.7/Pr/sub 0.3/(Fe/sub 1-x/Co/sub x/)/sub 1.85/ predominantly consists of the MgCu/sub 2/-type cubic Laves phase, with a small amount of PuNi/sub 3/-type and rare-earth rich phases. The amount of (Tb,Dy,Pr)(Fe,Co)/sub 2/ in (Tb/sub 0.7/Dy/sub 0.3/)/sub 0.7/Pr/sub 0.3/(Fe/sub 1-x/Co/sub x/)/sub 1.85/ increases with increasing Co concentration and becomes almost the only phase present when x/spl ges/0.4. The lattice parameter decreases approximately linearly with increasing x from 0.7357 nm when x=0 to 0.7300 nm when x=0.6. The Curie temperature increases with increasing x from 640 K when x=0 to 677 K when x=0.3, and then decreases to 572 K when x=0.6. The saturation magnetization M/sub s/ exhibits a minimum near x=0.3. The polycrystalline magnetostriction of (Tb/sub 0.7/Dy/sub 0.3/)/sub 0.7/Pr/sub 0.3/(Fe/sub 1-x/Co/sub x/)/sub 1.85/ decreases monotonically with increasing x. There is no evidence of an anomalously large /spl lambda//sub 100/ value as is the case for other R(Fe/sub 1-x/Co/sub x/)/sub 2/ compounds.

Microstructure evolution of rare earth rich phase for rapidly-solidified TiAl based alloys

Much attention has been paid to rapid solidification (RS) of titanium alloys these years. As potential structure material used at elevated temperatures, TiAl based alloys have been given much emphasis in the last two decades. There are two main ways of improving the poor ductility at room temperatures (RT) of TiAl based alloys. They are alloying technique and RS technique. In recent years, rare earth elements have been added to titanium alloys so as to enhance the properties at elevated temperatures through the forming of dispersed second phase particles. Rare earth rich phases are either rare earth oxides (RE2O3) or some intermetallics formed by rare earth elements and other metals such as AlCe [1–4]. In the experiment, rare earth mixture was added to TiAl based alloys. The purpose of adding rare earth elements to TiAl based alloys is obtaining fine dispersed second phase microstructure through RS techniques in order to improve ductility (RT) and mechanical propertied at elevated temperatures. Main constituent of the rare earth mixture is cerium and lanthanum. Table I is chemical analysis result of the alloy used in the experiment. Induction skull melting (ISM) was used in obtaining as-cast TiAl-RE alloy ingot. Through published papers [5, 6] we know that oxygen content in the alloy obtained by ISM can be easily controlled at the scope of 0.04–0.07 wt%. In the experiment, specimens used in RS technology were cut from the as-cast ingot. Two RS techniques employed in the experiment are melt spinning (MS) technique and hammer-and-anvil (HA) technique. The MS device used is 5 T Melt Spinner which was made by Marco Corp. of U.S.A. Ribbons of 50–80 μm thick and 30– 50 mm long were obtained. The cooling rate of MS technique for TiAl based alloys is estimated at the scope of 102–104 T/s. In the HA process, TiAl based alloy cubic of 2 mm × 2 mm × 2 mm was put on the copper anvil, after being arc-melted, the hammer was driven by high-pressure gas and hit the anvil with very high speed, finally, a thin foil (30–100 μm thickness) was obtained. The cooling rates can reach as high as 107 T/s [7]. Gas pressure of protective argon in the furnace is 0.03–0.05 Mpa. Though scanning electron microscopy (SEM) and transmission electron microscopy (TEM) microstructures of as-cast specimens and RS specimens were observed.

Production of (Nd,MM) 2(Fe,Co,Ni) 14B-type sintered magnets using a binary powder blending technique

The anisotropic (Nd,MM) 2(Fe,Co,Ni) 14B-type magnets were produced using the binary alloy sintering method (MM denotes a Misch-metal). The composition of the master alloy was close to stoichiometric Nd 2Fe 14B compound, while that of the sintering aid was MM 38.2Co 46.4Ni 15.4. The microstructures of sintered magnets were investigated using a scanning electron microscope (SEM) equipped with an energy dispersive X-ray detector (EDX). X-ray diffractometry (XRD) was also employed for phase analysis. The optimum magnetic properties were obtained for a composition of Nd 11.9MM 2.9Fe 73.9Co 3.3Ni 1.1B 6.9 made of 90 wt.% master alloy and 10 wt.% sintering aid. The two main phases, i.e. the (RE) 2(TM) 14B phase as the matrix and the rare-earth rich phase as the grain boundary phase, were identified in the specimens. No Nd 1.1Fe 4B 4 boride phase was observed in these samples, while there was some Fe-rich phase in the specimens with lower amount of sintering aid. The permanent magnet properties of the specimen in the optimum condition were as follows: B r=11.2 kG, H cj=11.5 kOe and ( BH) max=32 MGOe with a Curie temperature of 351°C.

Structure, magnetic properties, and magnetostriction of Sm0.88Dy0.12(Fe1−xMx)2 (M=Mn, Al) (0⩽x⩽0.2)

The structure, magnetic properties, and magnetostriction of arc-melted polycrystalline Sm0.88Dy0.12(Fe1−xMx)2 (M=Mn, Al, 0⩽x⩽0.2) alloys were investigated by microscopy, x-ray diffraction, A.C. initial susceptibility, vibrating sample magnetometry, and standard strain gauge techniques. It was found that the matrix consists almost entirely of the MgCu2-type cubic Laves phase, with a small amount of rare-earth rich phase. The lattice parameter increases approximately linearly with increasing x from 0.740 to 0.750 nm for M=Al. For M=Mn the lattice parameter increases by 0.002 8 nm over the same increase in x. The Curie temperature decreases with increasing x from 670 to 410 and 550 K for M=Al and Mn, respectively. The spontaneous magnetostriction, λ111, and polycrystalline magnetostriction, λs, both decrease with increasing Al or Mn substitution for Fe from their initial values of −1452 and −1010 ppm, respectively. Premature disappearance of the splitting of the (440) line and changes in the shape of the magnetostriction curves suggest the substitution of Mn for Fe may produce a change in the anisotropy. No such effects were observed for M=Al. The saturation magnetization, MS, decreases approximately linearly with increasing substitution of Al or Mn. The changes in the magnetization and magnetostriction can be understood on the basis of a decreasing 3d sublattice moment.

Structure and magnetostriction of PrxDy1−xFe2 and Pr0.4Dy0.6(Fe1−yMy)2 alloys (M=Co, Ni)

The structure, Curie temperature and magnetostriction of PrxDy1−xFe2 (0⩽x⩽0.5) and Pr0.4Dy0.6(Fe1−yMy)2 (0⩽y⩽0.6) alloys (M=Co, Ni) have been investigated using optical microscopy, x-ray diffraction, ac initial susceptibility and standard strain gauge techniques. The matrix of homogenized PrxDy1−xFe2 alloys is a cubic Laves phase (Pr, Dy)Fe2 with MgCu2-type structure, with a small amount of second phase (Pr, Dy)Fe3 when x⩽0.2. The amount of (Pr, Dy)Fe3 phase increases with the increase of Pr content, and it becomes the main phase when x=0.4. When x=0.5, the matrix is found to be the (Pr, Dy)2Fe17 phase coexisting with a small amount of phases (Pr, Dy)Fe2, (Pr, Dy)Fe3 and rare-earth rich phases. For Pr0.4Dy0.6(Fe1−yCoy)2 alloys, the amount of (Pr, Dy)(Fe, Co)2 phase increases with increasing Co content and the phase (Pr, Dy)(Fe, Co)2 becomes the main phase when y=0.6. However, the substitution of Ni for Fe up to 60 at % Ni in Pr0.4Dy0.6Fe2 alloys does not favor the formation of the cubic Laves phase (Pr, Dy)(Fe, Ni)2. The lattice constant of PrxDy1−xFe2 alloys decreases with increasing x, whereas the Curie temperature Tc increases. The magnetostriction of PrxDy1−xFe2 alloys at room temperature exhibits a peak at x=0.3. The lattice constant of Dy0.6Pr0.4(Fe1−yCoy)2 alloys decreases slowly with increasing y; Tc shows a peak when y=0.45, and the room temperature magnetostriction becomes negative when x>0.45. The Curie temperature of Dy0.6Pr0.4(Fe1−yNiy)2 alloys decreses with the increase of Ni content. The room temperature magnetostriction of Dy0.6Pr0.4(Fe1−yNiy)2 also becomes negative when x>0.45.

Structure and Magnetic Properties of Melt-Spun Pr(Fe1-xCox)2 Alloys

Pr(Fe1-xCox)2 ribbons have been prepared by melt-spinning method. Their structure, magnetic properties and thermal stability are investigated. It is found that Pr(Fe,Co)2 cubic Laves phase can form only when x≥0.2 at wheel speed of 45m/s. For Pr(Fe0.6Co0.4)2 alloy, at wheel speed of 30m/s, the ribbon consistes of a mixture of Pr2(Fe,Co)17, Pr(Fe,Co)2 and rare earth-rich phase. Almost Pr(Fe0.6Co0.4)2 compound a small amount of amorphous phase is observed at wheel speed up to 45m/s. Pr(Fe0.6Co0.4)2 compound becomes unstable above 770℃. The resub-bonded Pr(Fe0.6Co0.4)2 nanocrystalline compound which is obtained at wheel speed of 40m/s combines high magnetostriction (λ∥-λ⊥=140×10-6), with significant magnetic coercivity, iHc=398kA/m.

Structure, Curie temperature, and magnetostriction of Sm1−xPrxFe2 alloys

The structure, Curie temperature, and magnetostriction of Sm1−xPrxFe2 alloys have been investigated using optical microscopy, x-ray diffraction, ac susceptibility, diffractometer step scanning and standard strain gauge techniques. It is found that the matrix of homogenized Sm1−xPrxFe2 alloys is the (Sm,Pr)Fe2 phase with MgCu2-type structure up to x=0.3, and the minor phases, both (Sm,Pr)Fe3, and rare earth-rich phase appear. The lattice parameter of the alloys increases slowly with increasing Pr content, but the Curie temperature goes the opposite way. The room temperature magnetostriction |λ∥−λ⊥| and easy direction magnetostriction |λ111| of Sm1−xPrxFe2 alloys exhibit a peak near x=0.1.

Structure and magnetostriction of R(Fe 0.95A1 0.05) x alloys (R = Dy 0.65Tb 0.25Pr 0.1, 1.8 ≤ x ≤ 1.95)

The structure and magnetostriction of preferentially oriented polycrystalline R( Fe 0.95 Al 0.05) x alloys (R = Dy 0.65Tb 0.25Pr 0.1, 1.80 ≤ x ≤ 1.95) prepared by Czochralski technique, have been investigated using optical and SEM metallography, X-ray diffraction, transmission electron microscopy and standard strain gauge techniques. The alloy contains some rare earth-rich phase when x ≤ 1.85, however (Dy,Tb,Pr)(Fe,Al) 3 particles emerges when x > 1.85, the amount of which increases with increasing x. X-ray diffraction analysis indicated that almost all crystallites grow along 〈110〉 direction when x ≤ 1.85 whereas part of crystallites grows along 〈112〉 direction when x ≥ 1.90. The magnetostriction along the growth direction of the samples has been assessed.

Direct observation of the nucleation of rare-earth-rich phase particles in rapidly solidified Ti-5Al-4Sn-2Zr-1Mo-0.25Si-1 Nd alloy

Ti-5Al~4Sn-2Zr-lMo-0.25Si-1Nd is a near-alpha high temperature titanium alloy offering excellent tensile strength and good creep resistance up to 550°C together with improved fatigue strength. Addition of 1 wt% Nd to Ti-A1-Sn-Zr-Mo-Si system results in a uniform distribution of incoherent dispersoids and improves high temperature properties [1]. There have been many studies of rare earth additions to titanium alloys [2-6], but most work placed emphasis on binary or ternary alloys: none systematically reported Nd addition to high temperature titanium alloys such as the Ti-A1-Sn-ZrMo-Si system. It is well known that rapid solidification prevents kinetic growth of crystals, and even causes the formation of amorphous phases. It is possible that the nucleation process of the rare-earth-rich phase particles in Ti-5Al-4Sn-2Zr-lMo-0.25Si-lNd alloy can be clearly observed. By using an arc-melting piston and anvil apparatus, rapidly solidified Ti-5A14Sn-2Zr-lMo-0.25Si-lNd alloy flakes, 30-60/~m thick, were obtained. The observation was performed in a top-entry JEOL 2000 EXII high resolution electron microscope (HREM) with a point-to-point resolution of 0.21 nm. The foils were prepared by electrothinning in a methanol, butylethanol and perchloric acid mixture cooled to about -30 °C. A HREM image of a typical rare-earth-rich phase particle and matrix is given in Fig. 1. It can be seen that the matrix is in a crystalline state, and the rareearth-rich phase particle is amorphous. When the matrix has solidified, the rare-earth-rich phase remains liquid within a certain temperature range. According to the diffraction pattern, the matrix is a Ti, and the lattice fringes in the matrix are from the (0 0 0 2) plane. The contrast between the matrix and particle is mainly due to the lattice fringes (or to the rows of well-defined bright spots). EDAX results show that the amorphous phase consists of Nd and Sn, and the particles can be denoted as the rareearth-rich phase. Region V of Fig. 1 shows that a partial crystallization of the rare-earth-rich phase particle appears at the boundary between matrix and rare-earth-rich phase particle. Some rare-earth-rich phase particles begin to nucleate after matrix crystallization. It is clear that the nucleation of the rare-earth-rich phase Figure l HREM image of partially nucleated rare-earth-rich phase particle in the as-quenched Ti-5A1-4Sn 2Zl~lMo-0.25SilNd alloy.

Defects in Terfenol-D crystals

The detailed microstructure of Terfenol-D (Tb 0.27Dy 0.73Fe 1.95) single crystals, prepared by the Czochralski method using an induction-heated cold crucible, and a vertical float zoned Terfenol-D (Tb 0.27Dy 0.73Fe 1.9) rod has been studied by transmission electron microscopy (TEM). It is found that the single crystal prepared with the Czochralski method has a primary growth direction within 8° close to 〈112〉 and contains many fine Widmanstatten precipitates (WSP) distributed homogeneously across the crystal. On the other hand, the vertical float zoned rod consists of a few misorientated grains with dendritic morphologies and some rare earth-rich phases. TEM observation reveals that both types of crystals contain numerous {111} stacking faults. The vertical float zoned crystal contains a higher density of defects and a few interior oxide particles. Twins have been observed in the vertical float zoned crystals. The Widmanstatten needle phases in the crystal grown with the Czochralski technique are enclosed in interfacial dislocation networks, which are considered significantly to deteriorate the magnetostrictive properties.