VALVES EXHAUST GAS
Appendix 2 Appendix 2
2. Phase diagrams and crystal structures 1. Phase diagrams
The phase diagrams of some R - F e - B systems have been reported by various authors (Bilonizhko and Kuzma 1972, Kuzma et al. 1977, Chaban et al. 1979, Stepanchikova and Kuzma 1980, Stadelmaier and Park 1981, Stadelmaier et al.
1983 and 1984, Matsuura et al. 1985b, Oesterreicher 1985). In fig. 1 is plotted the phase diagram of the N d - F e - B system and in fig. 2 a section at T - - 600°C in the above phase diagram (Oesterreicher 1985). For Nd2Fe14B a melting temperature T M ~--1450°C is evidenced (Stadelmaier et al. 1984, Stadelmaier and E1 Masry 1985). The ternary borides crystallize from the melt. At the composition Nd2Fe14B, the primary crystallization product is 6-Fe. Annealing at 1200°C causes Nd2F%4 B to recrystallize into much larger grains. It also causes the amount of iron, which is now 7-Fe, to decrease by coalescence into fewer particles. One may observe the presence of a three phase region R2Fe14B- RI+~Fe4B4-R , with a low melting eutectic suitable for liquid-phase sintering.
-2200
(..9
#-
0-1000
1 5 4 0
j ' ,
1174 i "
T 1 5 3 5
i,
~
1637N2F°2 =
Fe B Fe2 ~ ~ N d
B
Fig. 1. Phase diagram of N d - F e - B system.
PHYSICAL PROPERTIES OF R2Fej4B-BASED ALLOYS N d ~
75
Nd Fe 2
"Nd s Fe~3 Nd2Fe~7 / d~
Nd~-xFe4+y
Fe Fe2B 50%B
Fig. 2. Phase equilibria in N d - F e - B system at 600°C.
According to Oesterreicher (1985) the RFe4B 4 compounds appear to allow for a slightly higher R concentration as compared to the stoichiometric one. Matsuura et al. (1985b) reported the existence of the RI+~Fe7B 6 composition. Crys- tallographic studies show the presence of the RI+~Fe4B 4 composition (R = Ce, Pr, Nd, Sm, Gd and Tb) with 0.11 (Pr)~< e ~<0.15 (Tb) (Bezinge et al. 1985, Braun et al. 1982, Givord et al. 1985a, b, Givord et al. 1986a, b). For R = Pr, Nd and Gd the structure is of the Pccn space group, while for R = Ce and Sm, the structure can be described by the space group P42/m. In addition, weak super- structure reflections were observed (Bezinge et al. 1985, Givord et al. 1985b, 1986a). These correspond to a periodic modulation of the Fe tetrahedra induced by interactions with the R sublattice in order to increase the shortest F e - R interatomic distances. Chang and Qiang (1986) reported that the tetragonal boron rich phase is one-dimensional.
An eutectic composition with low melting temperature is also shown. After Stadelmaier et al. (1984), the eutectic constituent found in the magnet alloys have 70 at% Nd with little boron.
The experimental study of the liquidus projection in the phase diagram of N d - F e - B ternary system has been also carried out (Matsuura et al. 1985b). In the Nd-poor region it is found that two phases, Fe2B and T 2 (Ndl+~Fe4B4 or Nd2Fe7B6) coexist with the liquid. Two monovariant curves L~,~Fe + Fe2B and L,~-F%B + T2, join and form the transition reaction, L + Fe2B ~ - T 2 + Fe. The end point of the liquid in the Nd-poor region is the ternary eutectic point
76 E. B U R Z O and H.R. KIRCHMAYR
L ~ - ~ - - F e + N d 2 F e 1 4 B ( T 1 ) + T 2. In the Nd-rich region T1 phase, T 2 phase and rich-Nd solid solutions solidify from the liquid at the end point of the liquid.
The phase diagram of as-cast buttons and melt-spun ribbons with the composi- tion Laa5FeyyB 8 and La16Fev6B 8 was investigated by Hadjipanayis et al. (1985c).
The crystallization and stability of the tetragonal LazF%aB phase was analyzed by Stadelmaier et al. (1985b). It has been suggested that in the P42/mnm-type structure boron can be replaced to some extent by carbon (Stadelmaier and E1 Masry 1985, Stadelmaier and Park 1981). An iron-rich phase having the composi- tion Gd3Fe20C was initially reported in the G d - F e - C system (Stadelmaier and Park 1981). The study of the G d - F e - C phase diagram allows one to identify this phase as Gd2Fe14 B (Stadelmaier and E1 Masry 1985). No such carbide was found in the C e - F e - C system (Park et al. 1982). This phase has been observed in the R - F e - C system with R = Dy or Er (Pedziwiatr et al. 1986c).
2.2. Crystal structure of R2Fe14B compounds
The crystal structure of Nd2Fe14B compound has been determined by neutron diffraction study (Herbst et al. 1984) and by X-ray measurements (Givord et al.
1984a, Shoemaker et al. 1984). This phase is tetragonal belonging to the space group P42/mnm (fig. 3).
The crystalline cell is constructed from four Nd2Fel4B units, 68 atoms, respec- tively. This structure is formed also with other lanthanides, yttrium and thorium.
The R ions occupy two distinct crystallographic sites, the iron ions share six non-equivalent sites, while B is located on one site. The positions of the atoms in lattice are given in table 1 (Herbst et al. 1984).
The above structure is related to the sigma phase (Bergman and Shoemaker 1954) and may be described in terms of nets perpendicular to z-axis. There are triangular nets formed by Ndl(4f), Nd2(4g), Fel(4e ) and B atoms in the mirror planes at z = 0 and 1/2. Between each pair of adjacent planes, there are sandwiched two puckered sigma-phase main layer-type nets (z---1/8 and 3/8) formed by Fe2(4c), Fe3(8jl), F%(16kl) and Fe6(16k2) atoms and one sigma-
TABLE 1
Atomic sites and coordinates x, y, z (in units of the lattice constants) for Nd2Fe14B obtained from analysis
of room temperature neutron diffraction data.
Atom Site x y z
Nd I 4f 0.2679(5) 0.2679(5) 0 Nd 2 4g 0.1403(4) 0.1403(4) 0
Fe 4c 0 0.5000 0
Fe 4e 0.5000 0.5000 0.1139(5)
Fe 8jl 0.0979(3) 0.0979(3) 0.2045(2) Fe 8j2 0.3167(3) 0.3167(3) 0.2464(3) Fe 16k I 0.2234(3) 0.5673(3) 0.1274(2) Fe 16k 2 0.0375(3) 0.3698(3) 0.1758(2) B 4g 0.3711(9) 0,37t1(9) 0
PHYSICAL PROPERTIES OF R2F%4B-BASED ALLOYS 77
Rig}
~ F e ( c l O F e { e / I D Fe(j.~l(~ Fe(j 2 } ~ F e (kl) ~ F e (k 2 ) @ B (g)
Fig. 3. The crystal structure of R2FeI~B compounds.
Fig. 4. The schematic representation of iron layers and ~r-phase showing the anti- prisms occupied by Fe4(8j2 ) atoms.
78 E. BURZO and H.R. KIRCHMAYR
TABI.E 2
The interatomic distances in Nd2Fe~4B Site Interatomic distances ( ~ )
Fel(4e ) 2 Nd2(4g ) 3.192 2 Fe3(8j~) 2.491
2 B 2.095 2 Fe(8j2 ) 2.754
1 Fe~(4e) 2.826 4 Fes(16kl) 2.496 Fe2(4c ) 2 Ndl(4f ) 3.382 4 F%(16k~ ) 2.573 2 Nd~(4g) 3.118 4 Fe6(16k2) 2.492 Fe3(8j, ) 2 Ndl(4f ) 3.306 1 Fe4(8jz ) 2.784 1 Ndz(4g ) 3.296 2 Fes(16kl) 2.587 1 Fel(4e ) 2.491 2 Fe6(16k2) 2.396 1 Fe3 (4f 1 ) 2.433
2 Fe~(8j2 ) 2.633
Fe4(8jz ) 1 Nd~(4f) 3.143 2 F%(16kl) 2.748 1 Nd2(4g ) 3.049 2 Fes(16k ~ ) 2.734 1 Fe~(4e) 2.754 2 Fe6(16k~) 2.640 2 Fe3(8j~ ) 2.633 2 Fe6(16k2) 2.662 1 Fe3(8jl ) 2.784
Fes(16k~) 1 Ndl(4f ) 3.066 1 Fe4(8j2 ) 2.734 1 Nd2(4g ) 3.060 1 Fe~(8j~) 2.748
1 B 2.096 1 Fes(16k~) 2.592
1 Fe~(4e) 2.496 1 Fe6(16k2) 2.527 1 Fe2(4c ) 2.573 1 F%(16k2) 2.536 1 F%(8j~) 2.587 1 Fe6(16k2) 2.462 Fe6(16k2) 1 Nda(4f) 3.279 1 F%(16kl) 2.527 1 Nd2(4g ) 3.069 1 F%(16kl) 2.462 1 Fe2(4c ) 2.492 1 Fes(16kl) 2.536 1 Fe3(Sj~ ) 2.396 1 Fe6(16k~) 2.542 1 Fe3(Sj2 ) 2.662 2 F%(16k2) 2.549 1 Fe4(8j2 ) 2.640
phase subsidiary layer-type net (z = 1/4) formed by Fe4(8j2 ) atoms (fig. 4).
Similarly, as in CaCus-type structure, the R atoms occupy the center of a hexagonal prism. The B atoms are at the center of a trigonal prism of iron atoms, occupying the lattice sites similar to the 2c-type in CaCu 5 structure.
The coordination of the atoms and also the distances between them, in Nd2Fe~4B, are given in table 2 (Givord et al. 1984a). In these systems short distances (d < 2.50 A) between some iron sites appear similar to those found in R2Fe~7-type compounds (Givord and Lemaire 1974).
2.3. Lattice parameters
of R2Fe14B
compoundsA survey of the composition dependence of the lattice constants of R2Fe~4B compounds was given already by Livingston (1985b). In fig. 5 we plotted the up-to-date reported values (Sagawa et al. 1984a, Givord et al. 1984a, Shoemaker et al. 1984, Abache and Oesterreicher 1985, Sinnema et al. 1984, Yamamoto et al. 1984, Stadelmaier et al. 1985a, Andreev et al. 1985, Burzo 1984, Hirosawa et
P H Y S I C A L P R O P E R T I E S O F R2Fe~4B-BASED A L L O Y S 79
12.z~
12.3
, , <
12.2
12.1 c
~12.0 c- o o
• - 11.9
8 . 8 o <
8.7 8.6
- 8 ~ * o
"~ u []
O
_ + 1:1
4-
•
+ ¢ 4
I t 'I ] I _ I t _ I I I __L I ~ I I I I
Y La Ce Pr Nd Pm Srn Eu Gd Tb Dy Ho Er Tm Yb Lu Fig. 5. T h e composition d e p e n d e n c e of the lattice parameters: O A b a c h e and Oesterreicher (1985),
* Sinnema et al. (1984), [] Sagawa et al. (1984a, b), • Yamamoto et al. (1984), A Shoemaker et al.
(1984), V Stadelmaier et al. (1985b), O A n d r e e v et al. (1985), O Burzo (1984), ~ Givord et al.
(1984a), + Hirosawa et al. (1985a), [] Deryagin et al. (1984).
al. 1985a, Deryagin et al. 1984). T h e s e decrease as the lanthanide atomic n u m b e r increases, following the lanthanide contraction. Such a behaviour is m o r e evident in the c direction. According to H e r b s t et al. (1985) these features are connected with the structural stability of the trigonal prisms f o r m e d by each b o r o n atom and its six nearest iron neighbours. T h e anomalously small lattice parameters of Ce2Fe14B c o m p o u n d are associated with the tetravalent state of cerium.
The crystal structure and lattice p a r a m e t e r s of RzFet4_xTxB alloys with T = Co, Si, A1, Cu, Mn, Cr, etc., were also studied. T h e R2Fe14_xSixB compounds (R = Pr, Nd, E r and Y) crystallize in a tetragonal structure of P42/mnm-type up to a silicon content x = 2. T h e lattice parameters decrease as the Si content increases (Pedziwiatr et al. 1987a, b). T h e RzFe14_xCoxB compounds with R = Pr, Nd, Gd, and Y have a tetragonal structure for the entire composition range (Sagawa et al. 1984b, Arai and Shibata 1985, Burzo et al. 1985a, b, Matsuura et al. 1985a, Pedziwiatr et al. 1986b, L ' H e r i t i e r and Fruchart 1985, Fuerst et al.
80 E. B U R Z O and H.R. K I R C H M A Y R
1986). The lattice parameters decrease with increasing the cobalt content. T h e R2Fe14B-type structure was preserved when replacing iron by aluminium (Burzo et al. 1986a, Yang et al. 1986a). Single phase compounds were obtained when substituting iron by ruthenium up to x = 2.0 (Pedziwiatr et al. 1986a, Ku et al.
1986). In Y2Fe~4 xCuxB alloys a single phase is observed up to a c o p p e r content x = 1.5 (Pedziwiatr et al. 1987a). Diffraction studies on Er2Fe~4_xTxB with T = Mn or Co show the presence of a tetragonal structure for T = Mn up to x = 9 and up to x = 5 for Co (Meisner and Fuerst 1986). T h e lattice p a r a m e t e r s of Er2Fe~4 xMnxB exhibit a minimum between x = 2 and x = 3. T h e lattice parame- ters of N d l s ( F e 1 xTx')7sB with x < 0.5 and T' = Mn or Cr increase with raising the Mn or Cr content (Yang et al. 1985a, b, H o et al. 1985). T h e structure of Gd2Fe~4_xMnxB ( L ' H e r i t i e r and Fruchart 1985) and Nd2Co2Fe~2 xMnxB (Jur- czyk and Wallace 1986) was also analysed. Maocai et al. (1985b) replaced Fe in Nd~5(Fe~_xT"x)77B 8 by Cr, Zr, Ti, Ni, Mn and Cu. No data on the structure of the above alloys were reported.
2.4.
Preparation of the alloys
In the following we present some preparation methods of R - F e - B alloys. The steps involved in manufacturing p e r m a n e n t magnets will be surveyed in section 4.1. Both the R2Fel4B magnetic phase as well as alloys having a composition slightly different from the above are obtained. For manufacturing p e r m a n e n t magnets, alloys with an approximate composition R~sFev7B 8 are used. The alloys are multiphases and always contain in addition to the R2Fe14B hard magnetic phase a low-melting eutectic phase with - 7 0 at% Nd as well as R 1 +~Fe4B 4 phase.
T h e most c o m m o n m e t h o d to produce R - F e - B alloys consists in melting the respective metals in an inert atmosphere, in induction furnace (Sagawa et al.
1984a) or in arc furnace ( T o k u n a g a and H a r a d a 1985). Sometimes, the melting is p e r f o r m e d in vacuum. T h e melting in vacuum has no advantage as c o m p a r e d to that in inert gas, since t h e r e is a loss of weight of some c o m p o n e n t s by volatilization. The Nd may be p r e p a r e d in the conventional way by electrolysis of the chloride, or thermal reduction of the fluoride by using calcium as the reductant. Variations include preparing a low melting Nd-rich N d - F e master alloy by the above two processes. T h e use of this master alloy is interesting technically as well as commercially.
Stadelmaier et al. (1985a) p r e p a r e d R - F e - B alloys by the reaction of elemental Fe and Nd p o w d e r in combination with a master alloy of FezB-type along a ternary diffusion path.
R2Fe~4B single crystals were also produced. Givord et al. (1984a) obtained crystals by the Czochralski m e t h o d from a levitation melt in a so-called cold crucible. K o o n et at. (1986a) p r e p a r e d R - F e - B single crystals using the Czoch- ralski m e t h o d and arcs to keep the charge molten. Typical growth rates were 5 m m / h . Hiroyoshi et at. (1985) have grown R2FeI4B single crystals in an infrared imaging furnace filled with argon by floating-zone melting technique.
P H Y S I C A L P R O P E R T I E S O F R2FeI4B-BASED A L L O Y S 81
After cooling the samples are annealed at temperatures between 900 and 1100°C in an inert gas atmosphere or in vacuum.
The R - F e - B alloys may be prepared by reducing with calcium, at high temperature, the respective rare-earth oxides, in the presence of a mixture of transition metals in powder form plus some transition metal oxides (Herget 1985a). Two versions of calciothermic processes for making rare-earth alloys were proposed: the reduction-diffusion ( R - D ) process (Cech 1974, McFarland 1973) and the co-reduction process (Herget and Domazer 1975, Domazer and Strnat 1976). The last one consists of a simultaneous reduction of rare-earth and transition metal oxides in the presence of transition metal powder, followed by diffusional alloys formation. The process may be divided into two steps:
~Nd203 q- 72Fe* + 340 Fe40B°60 + ~s Ca 1200°C NdlsFev7B8 + 2- CaO 45
*iron metal powder; o commercial powdered ferroboron, Ca ° + 2CaO + 2H20 RT) 2Ca(OH)2
o excess calcium metal from reduction.
Since the eutectic phase is a region of heavy corrosion, a leaching procedure has been developed (Herget 1985a) that enables us to remove the excess calcium and the CaO by-product in spite of the corrosive nature of N d - F e - B alloys.
Another procedure used to obtain R - F e - B alloys is the so called melt-spinning method. Thus, the amorphous precursor or the fine-grained metastable structure are produced directly. Such a method involves melting of the alloy or constituent elements in quartz tube. The melt under argon pressure is sprayed through an orifice in a quartz tube into a rotating water-cooled copper wheel or disc. The quench rate varies by changing the substrate velocity. Cooling rates in excess of 106K s -1 are produced (Ormerod 1984, 1985). A fine grained metastable struc- ture could be obtained by annealing treatment.
The method of liquid dynamic compaction (LDC) (Chin et al. 1986, Tanigawa et al. 1986) is based on the process of gas atomization (Anand et al. 1980). In gas atomization a stream of molten alloy is broken into a spray of fine particles by a jet of high-velocity gas and the rapidly solidified particles are collected. In LDC, a cooled substrate is placed beneath the atomization core at a distance such that most of the sprayed droplets are partially solidified. The rapidly solidified alloy builds up on the substrate at controllable rates, which can easily exceed 1 cm/min.
Rapid solidification is made possible by the supercooling of the high-velocity atomized particles and the good thermal contact with a water-cooled copper substrate.
82 E. B U R Z O and H . R . K I R C H M A Y R