The ErY, TbHo, TbEr, DyHo, DyEr and HoEr phase systems

Journul 41 rhe Le.5,5-Conwwn Metals, (‘ Elsevier Sequoia S.A.. Lausanne

THE Er-Y,

Tb-Ho,

F. H. SPEDDING.

Tb-per, DyyHo,

B. SANDEN

August

Dy-Er

AND

Ho-Er

PHASE

SYSTEMS

Prliwrsiry.

Andes. fwu

and B. J. BEAUDR!

Anwv Ltrhorufo,!,~~~~:SAEC and Departrnenr (L’XA.) (Received

1

31 (1973) l- 13 Printed in The Netherlands

of’ Metu/lurg~,

low

Srate

50010

7. 1972)

SUMMARY

The ErrY, Tb-Ho, TbbEr, Dy-Ho, DyyEr and Ho-Er systems have been studied by thermal, X-ray and metallographic methods. It is known that these properties can be influenced significantly by the presence of non-metallic impurities. Therefore, these measurements were made on very pure metals prepared in this Laboratory and the alloys were carefully analyzed for all possible impurities. These systems are characterized by complete mutual solubility and within our experimental errors, no enrichment of one rare earth over the other was gbserved during solidification or transformation. The solidus and liquidus lines, therefore, fell on top of one another over the entire composition range. The addition of Er to Y or Y to Er lowers the melting point which reaches a minimum at 53 at.S(, Er. Alloys containing less than 53 at.% Er solidify to a b.c.c. lattice which transforms at lower temperatures to a h.c.p. lattice. Alloys with greater than 53 at.% Er crystallize directly to the h.c.p. solid solution. For the Tb-Ho, Tb-Er, Dy-Ho and Dy-Er systems, the melting point appears to decrease linearly with increasing at.“/; of the lower atomic numbered rare earth. However, when the atomic number of the components are two or more apart, there may be a very slight negative deviation from linearity which reaches a maximum at the composition where the b.c.c. first starts to form from the melt. This is especially evident in the DyyEr system in which the negative deviation could be interpreted as a slight change in slope occurring at the composition where the b.c.c. first starts to form from the melt. The b.c.c. starts to form at 10 and 34 at.% Tb in the Tb-Ho and Tb-Er systems and at 25 and 50 at.;< Dy in the DyyHo and Dy-Er systems, respectively. The transformation temperature is lowered linearly from these points with increasing composition of the lighter rare earth to the transition temperature of the pure solute. The Ho-Er system has a linear variation of melting points with composition. The a and c lattice parameters vary linearly with composition within our experimental errors for the binary systems of adjacent elements. By the time the elements are four apart in the series, Tb--Er, a distinct deviation from linearity occurs in the plot of lattice constant IWSUS composition which is positive for the a axis and negative for the c axis.

2

F. H. SPEDDING,

B. SANDEN, B. J. BEAUDRY

INTRODUCTION

The rare-earth metals which crystallize in the hexagonal close-pack ABAB structure at room temperature form solid solution binary alloys where the one component is completely soluble in the other over the entire composition range. A number of these systems determined with less pure materials have been reported’-lo. Since the room temperature solid solutions formed between lanthanides of the ABAB structure are nearly ideal, the high temperature solutions should exhibit a linear variation of the melting and transformation temperatures for these systems. The rare-earth metals scandium and yttrium, which are not lanthanides, also form solid solutions when alloyed with each other or the ABAB lanthanides. However, a plot of melting points versus composition for these systems sometimes shows a deviation from linearity, as might be expected from their difference in electronic structure where fewer shells are completed. As the size disparity of the two components increases, this deviation from linearity becomes more marked2x3. There seems to be a considerable discrepancy in the literature concerning the details of the intra-rare-earth binary systems. For example, the Y-Gd’ and Y-Ho systems4 were shown to have a linear relationship while the Y-Tb4 and Dy-Ho systems7 had slightly negative deviations in their melting and transition temperatures versus composition. In addition, the Dy-Y system4 was reported to have minima nearly 100 degrees below the expected melting and transition temperatures. This behavior would not have been expected from empirical rules. A study of the Y-Er system’ showed a peritectic horizontal at 1500°C with a large separation of the solidus and liquidus. Another investigation of the Y-Er system6 reported a continuous series of hexagonal solid solutions at low temperatures. However, the effect of erbium on the transition temperature of yttrium was not reported. Holmium was reported to have an anomaly in the resistivity vs. temperature curve at 1415°C” although no thermal arrest was noted in the same study on a separate sample. The authors” pointed out that impurities could have caused the anomaly in resistivity. It is well known that the rare-earth metals are powerful getters for the non-metallic impurities oxygen, nitrogen, carbon, and hydrogen and that these impurities can influence the experimental values of many properties. Unless extraordinary care with careful analytical control is maintained while preparing and handling these metals, the rare-earth metals are certain to contain an appreciable atomic percent of these non-metallic impurities and it is these impurities which we believe are responsible for the discrepancies of the data reported in the literature. Recently, this Laboratory has succeeded in preparing rare-earth metals where the nonmetallic impurities are very low, and the present study was undertaken to establish precisely the melting and transition temperatures of these rare-earth metals and their binary alloys using well characterized samples and specially designed thermal analysis equipment. The transformation temperature in various holmium alloys was of interest to determine if a transformation in holmium could be predicted by extrapolating to 100% holmium. The wide separation of the solidus and liquidus enrichreported in some of the previous studies4-7, which indicated a considerable ment of one rare-earth component over the other when the alloy solidified or

Er-Y, Tb- Ho. Tb per. Dy-Ho.

Dy-Er. Ho-Er SYSTEMS

3

transformed in crystal structure, seemed contrary to the general experience encountered in trying to separate the rare earths, since in all other cases such operations had to be repeated many times by fractionation methods to achieve a reasonable enrichment. We were therefore interested to see if such separations were real or due to the presence of impurities. Since a method has recently been established for the proper handling of rare-earth metals and alloys for accurate lattice parameter determinations”. a precise determination of the variation of the lattice parameters with composition was of interest to determine the departure, if any, of u0 and c0 wrsus composition from a straight line. EXPERIMENTAL

The metals employed in this investigation were prepared in this Laboratory using methods described previously’ 3. 14. In these methods, highly purified fluoride is reduced with calcium relatively free of non-metallic elements in a tantalum crucible. The reduced metals are purified by vacuum casting followed by distillation for yttrium and terbium and sublimation for dysprosium, holmium and erbium. The analyses of the metals used are given in Table 1 in parts per million atomic as well as weight ppm. The oxygen content of the yttrium and terbium is somewhat higher than that for the other metals because the distillation or sublimation processes are less efficient in separating yttrium and terbium from their oxide impurity than for the other members of the rare-earth series. In this part of the rare-earth series. the vapor pressure of TbO and YO are very close to the vapor pressures of the pure metals. Thermal methods

The accurate determination of melting points and transition temperatures of the rare-earth metals and alloys is dependent on many variables. Unless great care is taken in the design of the experimental equipment and in the method of determining these constants, considerable errors can result in the values reported. The thermocouples must be of known calibration and care must be taken to prevent their contamination by metallic vapors during a run. At the high temperatures encountered in this study the breakdown of the insulating properties of ceramic insulators frequently introduces electrical feedback which can raise or lower the e.m.f. produced by the thermocouple. An isothermal arrest during melting or transformation is usually an indication of high purity. However, a non-isothermal arrest can also be caused by poor thermal contact between the sample and thermocouple, as well as thermal gradients within the furnace. In our studies these variables were taken into consideration in the design of the equipment and the experimental runs. The furnace used was a tantalum tube resistance furnace which had a heat zone 6.4 cm diam. x 22.9 cm long. Radiation shielding at each end of the furnace helped to produce a 14.0 cm long constant temperature zone in the center. Thermal gradients over this length of the heater were further reduced by using a heavywalled molybdenum sample chamber, as shown in Fig. 1. The sample crucible and a crucible of the same dimensions filled with lutetium were placed in the holes in the molybdenum as shown. In a differential thermal apparatus, it is important that the

4 TABLE TYPICAL ERBIUM

F. H. SPEDDING,

B. SANDEN,

B. J. BEAUDRY

I ANALYSIS* OF THE YTTRIUM, TERBIUM, DYSPROSIUM, IN PPM ATOMIC (PPM BY WEIGHT IN PARENTHESES)

HOLMIUM

AND

Metal Impurity

Yttrium

o**

1251

N** H** F*** C*** Mg Al Si SC Ca Cr Fe Ni CU Y Ta La Ce Pr Nd Sm Gd Tb DY Ho Er Tm Yb LU

8:: 421 333 4 30 19 49 4 10 13 15 14 3 0.6 13 0.3 1 0.2 2 0.6 0.6 1 0.5 0.1 0.1 1

(225) (r:; (90) (45)

Terbium

Dysprosium

Holmium

Erbium

1340 (135) 46 (4) 315 ii; 33 66 (5) 0.7 6 11 18 20 6 57 8 20 11 53 7 3 2 4 9 2

581 ‘?ii

556 59 816 26 21 2 2 59 22 0.8 6 30 3 20 I 5 4 0.1 0.4 3 0.1 2 1 6

368 12 331 26

2 10 3 10 6 1

46 161 1;;

(1)

(r:;

1 12 6 4 4 6 41 11 20 4

1 4 2 2 11 1 3 5 20 15 0.3 0.2 0.3

20 0.5 0.5 0.2

;:; (2)

;:

97 2 12 11 11 0.8 13 12 3 145 28 60 2 1 2 2 0.1 4 0.5 15 15 4 0.7 1

All analyses unless otherwise specified were made by a spark-source mass spectrometer method should be accurate to a factor of 2 or 3 at the concentrations repo’rted. Elements not listed determined to be present in amounts less than I ppm atomic in all five metals. **0 , N > and H were determined by vacuum fusion and are believed to be reliable to +20% at levels. ***F was determined by the distillation and spectrophotometric determination of fluosilicic acid carbon by a combustion-chromatographic technique. l

(35)

and were these and

reference sample should have about the same heat capacity as the sample under investigation. Then, except when a transformation is taking place, the two samples will heat and cool at the same rate. For this reason lutetium was chosen as the reference sample. The crucibles were supported by the thermocouples and were not allowed to touch the molybdenum block, which prevented electrical feedback. After initially heating the system under vacuum to the maximum temperature, about 20 mm of helium gas were bled in to further insure thermal equilibrium between the sample and thermocouple.

Er- Y, Tb-Ho.

Tb- Er. Dy-Ho,

Dy-Er,

5

Ho Er SYSTEMS

MOLYBDENUM

THERMOCOUPL IS

I”4

I-

SCALE Fig.

biew showing the position and size of the sample block used to minimize thermal gradients.

1. A cut-away

molybdenum

and

reference

crucible

in the

Tantalum crucibles 1.27 cm diam. x 3.81 cm long with thermocouple wells 0.32 cm x 1.59 cm long in the bottom of the crucible were used. The samples of the pure metals were prepared in the form of cylinders by arc melting on a watercooled copper mold. A hole 0.32 cm x 1.59 cm was drilled into the end of the samples to allow them to be placed over the thermocouple wells in the bottom of the tantalum crucibles. To prevent contamination of the thermocouple by rare-earth metal vapors, a lid was welded on the tantalum crucible after evacuation and backfilling with helium in a welding chamber, as described by Miller et al.“. Data taken on the transformation temperature while heating and cooling the metal before melting were the same as after melting within the one degree precision of our instrument. The initial melting point was also the same as the melting point observed on repeated cycles indicating little or no effect of tantalum on the melting point and transition temperatures of all the metals studied. The data in Table I for the tantalum content is for the starting metals and does not show the final tantalum content. Dennison et ul. 16, l7 have studied the solubility of tantalum in the rare-earth metals and have shown the solubility ranges from 0.15 to 0.44 at.“,, at the melting point of the metals reported here. Therefore, the specimens undoubtedly contained some tantalum after they had been cycled several times through the melting point but, as pointed out above, any solubility of tantalum had a negligible effec.t on the melting and transition temperatures. To prepare the alloys, weighed portions of the metals were placed in a Ta crucible which was then evacuated, backfilled with helium and a lid welded in place. The crucible and contents were placed in the DTA apparatus and heated to 50

F. H. SPEDDING,

6

B. SANDEN, B. J. BEAUDRY

above the highest melting point. Preparing the samples by induction melting in the tantalum crucible or by arc melting on a water-cooled copper hearth had no noticeable effect on the thermal analysis results. The thermocouples used were PttPt 13”/,Rh with high-purity (99.85%) Al,O, insulators. After each run above 1200°C a new portion of thermocouple wire was taken into the hot zone to prevent accumulative contamination of the thermocouple bead from affecting the readings. Each new thermocouple was checked at the copper point for accuracy and the roll of wire was checked by having a calibration curve made by the National Bureau of Standards on a thermocouple made from the roll. The certification indicated an accuracy of k 1” up to 1450°C. Below 1100°C the calibration was obtained by comparison with a standard thermocouple while above 1100°C the calibration was obtained by extrapolation as discussed in NBS Circular 5901*. The temperatures reported which are above 1450°C were read from standard were converted to conform e.m.f. versus temperature tables’ 9. The final temperatures to the Practical Temperature Scale of 1968 ” . The precision of + 1.5”C obtained on points above 1450°C with different sections of thermocouple wire and new samples (in the case of the pure metals) leads us to believe an accuracy of +2”C was obtained at these higher temperatures. The precision on repeated heating and cooling cycles on each sample was better than one degree. The e.m.f.s were recorded continuously on a two pen recorder. The differential temperature was recorded on a - 1 to + 1 mV range while the timeetemperature pen had a full scale of 5 mV. To read higher e.m.f.s, a negative potential was manually added in increments of 5 mV to extend the range of the recorder. During a thermal arrest, the e.m.f.s were measured with a Leeds and Northrup K-3 potentiometer to obtain greater accuracy of measurement. During this investigation it was noted that frequently the thermocouple bead adhered to the tantalum thermocouple well after runs to 1525°C or higher, causing thermocouple breakage when the crucible was removed. Examination of one of these thermocouple assemblies showed a deposit around the bead and a deterioration of the bead and the wires just below the bead. The deposit was shown by emission spectroscopy to be rich in aluminum. While at first sight most reactions would appear to be energetically unfavorable, apparently the Al,O, insulator was reacting with the tantalum to form aluminum or a volatile aluminum compound which condensed around the thermocouple bead forming a low melting platinumaluminum-tantalum weld. The problem was solved by wrapping a 0.025 mm foil of platinum around the Al,O, insulator, isolating it from the tantalum. X-ray

methods

Samples to be studied by X-ray diffraction were prepared by co-melting the two metals on a water-cooled copper hearth in a nonconsumable electrode arc-melting furnace under an atmosphere of purified argon. The samples weighed between three and four grams. Typical weight losses during preparation were less than 0.01 g and were considered negligible. The as-arc-melted buttons were sandwiched in tantalum and cold rolled about 25% reduction in thickness. The sample was electropolished to clean the surface, wrapped loosely in tantalum, and heated at - 1 x lo-’ Torr to 550°C for 20 h. After cooling, each button was sampled for vacuum fusion analysis and metallographic examination. Two specimens 3 mm x 3 mm x 6 mm were cut

Er-Y. Tb-Ho.

Tb-Er.

Dye-Ho, Dy-Er.

7

Ho-per SYSTEMS

from the button and each prepared for X-ray diffraction by filing one end into a cylinder 2 mm long x 0.5 mm diam. The cold worked surface introduced by filing was removed by electropolishing in a 6:; perchloric acid in methanol solution at approximately - 70°C as described previously’ 2. The final diameter of the cylindrical portion of the sample was about 0.3 mm. This portion was centered in the X-ray beam in a 114.6 mm diam. Debye-Scherrer camera. Diffraction patterns were taken using filtered Cu radiation for Ho, Er, Y and their alloys while filtered Co radiation was used for Tb, Dy and alloys rich in these two elements. The lattice parameters were obtained by the use of a computer program21 which applies a least square fit to the extrapolation of back reflection data using a NelsonRiley function. The metallographic samples were prepared for examination using the electropolishing method of Peterson and Hopkins 22. The samples were etched by holding them in the electropolishing solution for one minute without voltage applied. RESULTS

The phase systems shown in Figs. 2-7 were drawn from the results of thermal and X-ray studies. The thermal results are shown as single points for the melting temperatures and for the transition temperatures, because the start and stop

3.61

-

OQ 360.c o 3.59

-

y

Fig.

2.

10

20

The Y-Er

Fig. 3. The Tb-Ho

30

40

AT.

%

50

60

70

60

90

Er AT

Erbium

system and lattice parameters system and lattice parameters

wrs,u wr.us

composition. composition

%

Holmium

8

F. H. SPEDDING, B. SANDEN, B. J. BEAUDRY

~1500 E $1400 E 1356 % g I300 1289

SOLID

SOLUTION i

!

5.700

A = 0, 0 = 0,

3.620

I

4 “0. CJ zC,

5.660

3.600

-

5.650

-

5.640

ILQ3.580 .s d 3.560 IO

20

30

40

AT.

%

50

60

Erbium

70

a0

90

100 Er

o

IO

DY

20 30 40 50 &I. % Holmium

60

70

80

90

100 Ho

Fig. 4. The Tb--Er system and lattice parameters oers~s composition. Fig. 5. The

Dy--Ho

system and lattice parameters LWSUScomposition.

of the thermal arrests were within 1.5” on both the heating and cooling curves, the same as the pure metals. The composition of the solid and liquid, therefore, are the same within our detection limits. The addition of Er to Y and Y to Er lowers the melting point which reaches a minimum of 1507°C at 53 at.% Er. The transformation temperature of Y is raised by the addition of Er. The transformation temperature curve was extrapolated to the solidus with the intersection occurring at 53 at.% Er. (See Fig. 2.) The transformation temperature of Tb is raised linearly by the addition of Ho but at a greater slope than the melting temperatures. The point at which the two curves intersect was determined by extrapolating the transformation temperature curve to the solidus. The intersection occurs at 90 at.% Ho in Tb. (See Fig. 3.) This behavior confirms the absence of a high-temperature b.c.c. form in Ho which is relatively free from non-rare-earth impurities. The addition of Er to Tb raises the transformation temperature of Tb linearly with the curve intersecting the solidus at 66 at.% Er. (See Fig. 4.) The same behavior was noted in the Dy systems with Ho and Er with the transformation temperature curve intersecting the solidus at 75 at.% Ho in Dy (see Fig. 5) and 50 at.% Er in Dy (see Fig. 6), respectively. The Ho-Er system has a linear variation of melting points with composition. (See Fig. 7.) The melting points and transition temperatures for the pure metals which are shown in Figs. 2-7 and reported in Table II were determined in this study and were

Er- Y, Tb-Ho.

& 3

Tb-Er.

Dy-Ho,

Dy-Er.

HopEr

9

SYSTEMS

1450

t “w 9 1400 1412 P

1381

t

A’a.

-

o=co

-

5.620

- 5.640

5.610

- 5.620

4

5.60;;

5 590

0

IO

DY

20

30

40

50

AT

%

Erbum

Fig. 6. The DyyEr

60

70

80

90

100 Er

0

HO

I I IO 20

system and lattice parameters

wrsu~ composition.

Fig. 7. The Ho Er system and lattice parameters

~ws~~.scomposition.

TABLE

I

I

30 AT

40 %

I

I

I

I

50 60 Erbwm

70

80

90

I

I 100 Er

II

MELTING

POINTS

AND TRANSITION

TEMPERATURES

OF THE METALS

STUDIED

Melt11

Terbium Dysprosium Holmium Erbium Yttrium

1356+2 1412k2 1474*2 lj29*3 1522*3

1289&2 13811-2 None observed None observed 1478+2

corrected to conform to the International Practical Temperature Scale of 196S2”. The agreement between the values given here and the best literature values is very good for Tb, Dy and Ho”. However, the melting point of Er is 5’ higher and the transformation temperature and melting point of Y about 5” lower than the best literature values”.23. Since our studies were done with carefully calibrated thermocouples on highly purified samples, we believe the results reported here are the most reliable reported thus far. The lattice parameters of each of the systems are plotted in the lower portions of Figs. 2-7. Since the exact lattice parameter cannot be read from these Figures

=I

10 TABLE

F. H. SPEDDING,

B. SANDEN,

B. J. BEAUDRY

III

MEASURED

AND CALCULATED

Composition (at.% Er)

LATTICE

PARAMETERS

Measured*

AND c/a RATIOS Calculated

c/a

a

c

r/a

5.7318*6 5.7139&2 5.6953 f 2 5.6779 f 2 5.6604 k 1 5.6409 k 1 5.6226 k 2 5.6038 k 1 5.5850+_3

1.5711 1.5710 1.5705 1.5705 1.5702 1.5699 1.5697 1.5693 1.5692

3.6482 3.637 1 3.6256 3.6148 3.6040 3.5924 3.58 14 3.5703 3.5592

5.7318 5.1136 5.6945 5.6768 5.6588 5.6398 5.6216 5.6034 5.5850

1.5711 1.5709 1.5706 1.5704 1.5702 1.5699 1.5697 1.5694 1.5692

Terbium-Holmium System** Tb 3.6055*4 9.9 3.6037+4 19.9 3.6008*3 29.9 3.598 1 * 1 39.9 3.5958+3 49.9 3.5929 k 3 59.9 3.5900* 1 69.9 3.5868*2 79.9 3.5838+2 90.0 3.5815+2 Ho 3.5778 f 2

5.6966 k 5.6897 k 5.6813+2 5.613 1+ 3.6646 k 5.6512 k 5.649 1 k 5.6415*2 5.6341 k 5.6268 + 5.6178*3

6 3

1.5800 1.5788 1.5778 1.5767 1.5753 1.5745 1.5736 1.5728 1.5721 1.5711 1.5702

3.6055 3.6028 3.6000 3.5912 3.5944 3.5917 3.5889 3.5861 3.5834 3.5806 3.5778

5.6966 5.6888 5.6809 5.6730 5.6652 5.6573 5.6494 5.6415 5.6336 5.6251 5.6178

1.5800 1.5790 1.5780 1.5770 1.5761 1.5751 1.5741 1.5731 1.5721 1.5710 1.5702

Terbium-Erbium Tb 11.3 17.4 26.3 38.8 46.1 58.6 67.6 80.3 87.0 Er

5.6966 k 5.6833 + 5.6768 + 5.6661+ 5.6524 k 5.6438+2 5.6299_+ 5.6200 + 5.6060 k 5.5996_+ 5.5850*

6 5 1 4 1 1 1 1 1 3

1.5800 1.5783 1.5776 1.5762 1.5747 1.5738 1.5724 1.5716 1.5706 1.5700 1.5692

3.6055 3.6002 3.5974 3.5933 3.5875 3.5842 3.5784 3.5742 3.5683 3.5652 3.5592

5.6966 5.6840 5.6712 5.6611 5.6533 5.6452 5.6312 5.6211 5.6070 5.5997 5.5850

1.5800 1.5788 1.578 1 1.5772 1.5758 1.5750 1.5136 1.5727 1.5713 1.5706 1.5692

5.650 1 k 4 5.6466 k 1 5.6432 f 1 5.6402 + 2 5.6368 k 5 5.6338 k 4 5.6298 + 1 5.6212 & 1 5.6226 k 1 5.6178*3

1.5732 1.5729 1.5725 1.5721 1.5717 1.5714 1.5710 1.5709 1.5706 1.5702

3.5915 3.5901 3.5886 3.5874 3.5859 3.5845 3.5830 3.5818 3.5797 3.5778

5.6501 5.6468 5.6433 5.6403 5.6368 5.6337 5.6301 5.6212 5.6222 5.6178

1.5732 1.5729 1.5726 1.5723 1.5720 1.5717 1.57 14 1.5711 1.5706 1.5702

a0 YttriumPErbium System** 0 3.6482 + 12.4 3.6312 k 25.4 3.6264+ 31.5 3.6154* 49.7 3.6050 k 62.7 3.5932 k 75.1 3.5820+2 87.5 3.5708 + Er 3.5592f2

2A 1 1 1 1 1 1

System*** 3.6055 f 4 3.6010*2 3.5984*5 3.5952 f 1 3.5896 k 3 3.5862 + 2 3.5805+2 3.5761& 2 3.5693 + 3 3.5665&3 3.5592 f 2

Dysprosium-Holmium System*** 3.5915*2 DY 10.2 3.5900*2 21.1 3.5886+ 1 30.3 3.5876+2. 41.0 3.5864 k 3 50.9 3.5851 k 1 62.0 3.5836+ 1 70.7 3.5823 + 2 86.2 3.5799* 1 Ho 3.5778 + 2

3 5 2 1 2 1

(Continued)

Er Y. Tb- Ho, Th Er. Dy-Ho, TABLE

DypEr.

Ho-Er

11

SYSTEMS

III /uwrd.)

Dj~\prosinm Ehmf Dy

Il.3 77 1 __.32.2 3X.5 47. I 51.h 6X.7 76.X XX.9 Er

Sj~.stem*** 3.5915+2 3.5X84_+? .3.5X50+4 3.5X22*2 3.%02+_2 3.5773*2 3.5739+ 3 3.5705+2 3.56X6&2 3.5632+2 3.559212

Holrnititli-EPhittrtl Sy.\retn*** 3.5778 * 2 Ho 3.5758 + 2 11.6 3.5742+2 20.9 3.5728 + 2 31.X 3X.1 3.5711+2 3.5692 * I 37.4 3.5674* 3 56.0 3.5657+_ 3 6X.0 3.5633 + 2 80.9 3.56172 I XX.4 3.5592+ 2 Er -. _~

5.650 I + 4 5.6427 + 4 .i.b350* 5 5.6284+4 5.6236&Z 5.6182+4 5.6119*4 5.6056i3 5.600X+3 5.5924*2 5.5850+3

5.6’17813 5.6142*2 5.61 12k I 5.6078 + 3 5.6058 *4 5.6026 & 3 5.59991_2 5.5965 t_ 3 5.5918t_2 5.5X96+2 SSXSO& 3

1.5732 1.5726 I.5718 1.5712 1.5708 1.5705 1S702 I.5700 I S695

1.5695 1S692

1.5702 1.5701 1.5699 1.5696 I .569X I .5697 I .5697 I S695 1Sb93 1.5694

1.5692

3.5915 3.5X7X 3.5833 3.581 I 3.5791 3.5763 3.5729 3.5693 3.5667 3.562X 3.5592

5.6501 5.6427 5.6356 5.6292 5.6250 5.6194 5.6126 5.6053 5.bOO I 5.5922 5.5X50

3.5778 3.5756 3.5739 3.5739 3.5707 3.5690 3.5674 3.565 I 3.5627 3.5613 3.5592 -.

5.6 I 7X 5.6140 5.6109 5.6074 5.6053 5.6022 5.5994 5.5955 5.5913 5.5X8X 5.5850 -.

a f 2 E 0.0002 * The lattice parameters of the pure metals are those reported in ref. 12 and the error limits maximum daiation from the average of 4 determinations. ** The limits of error given for the Er--Y and Th-Ho alloys are the standard deviations determination. *** The lattice parameters given for the Tb-Er. Dv Ho, Dv Er and Ho Er systems are the of two determinations. The limits of error for these systems are the deviations from the

1.5732 I.5727

1.5713 1.5719 lS7lb 15713 1.5709 I.5704 1.5701 I.5696 1.5692

1.570~ 1.570; I .57OO

I .5699 I .569X I .569? I .569b 1.5695 I .5604 I .5693 1.5692

arc the for one average average.

but only the trends shown, the lattice parameters measured as well as the lattice parameters calculated for ideal behavior are listed in Table III. As will be noted, for the binary systems of adjacent elements. (that is Dy-Ho and HoPEr) the variation of u and c with composition is essentially linear. When the components are four apart in the series. such as the Tb-Er system, a distinct deviation from linearity occurs in the plot of lattice constant ~‘ersus composition which is positive for the u axis and negative for the c’ axis. In the case of the Y.-.Er system, both the II and c lattice parameters have a slightly positive deviation from linearity which results in a linear (‘;iu ratio. DISCUSSION

The continuous series of sohd and liquid solutions which were observed in the present study were expected from previous studies made on these systems. The

12

F. H. SPEDDING,

B. SANDEN,

B. J. BEAUDRY

isothermal arrests observed for the melting and transformation temperatures show that there is no appreciable enrichment of one component over the other during these processes, which is what would be expected from the general experience encountered in trying to separate the rare earths. The lowering of the melting point of Y by Er and vice versa was not expected. However, the melting points and transition temperatures observed were as sharp as for the pure metals which indicates no appreciable enrichment occurs in this system either. The wide separation between the solidus and liquidus lines for these alloys, reported in some previous papers, we believe are not due to the enrichment of one rare earth over the other during freezing but to the presence of impurities, particularly non-metallic ones. We also believe that the pronounced minima in the melting point or transition temperature of the heavy lanthanide alloy series which have been reported, do not occur when the component rare-earth metals are pure, but may well result if considerable impurities are present. The TbbHo and Dy-Ho systems both show the absence of a b.c.c. form of holmium. That is, the transformation temperatures in these systems extrapolate to the solidus, not some temperature below the melting point of holmium. The composition at which the transition temperature curve intersects the solidus in the Tb-Er and Dy-Er systems also has an effective atomic number slightly below that of holmium. Whenever a b.c.c. phase has been reported in holmium at atmospheric pressure, we believe it resulted from the presence of impurities which lower the free energy of the b.c.c. phase with respect to the h.c.p. phase. For instance, the resistivity anomaly in Ho reported earlier was probably due to trace impurities, as suggested by Habermann and Daane’ ‘. Interstitial impurities, particularly hydrogen, can have a marked effect on the lattice parameters of the rare-earth metalsi2. Great care must be taken to prevent the introduction of hydrogen into the metal lattice during the preparation of samples for X-ray analysis. If impurities, such as hydrogen, are introduced or are present in the starting materials in varying amounts, the deviations of a and c versus composition from a linear relationship can be masked by their effect. In the present study great care was taken to use pure starting materials and to keep them pure. The alloys in these systems which simulate holmium (50 at.% Er in Dy and 66.6 at.% Er in Tb) have both a and c lattice parameters quite close to holmium metal. However, the alloys which simulate dysprosium (50 at.% Ho in Tb and 33.3 at.% Er in Tb) have a parameters similar to dysprosium while the c parameters TABLE LATTICE

IV PARAMETERS

OF ALLOYS

WHICH

SIMULATE

Sample composition

a0

co

cla

Ho 50 at.% Er in Dy 66.6 at.% Er in Tb

3.5778 3.5763 3.5766

5.6178 5.6163 5.6211

1.5702 1.5704 1.5717

DY 50 at.% Ho in Tb 33.3 at.% Er in Tb

3.5915 3.5929 3.5921

5.6501 5.6572 5.6587

1.5732 1.5745 1.5754

HOLMIUM

AND

DYSPROSIUM

Er-Y. Tb Ho, Tb-per. Dy--Ho.

Dy-Er,

Ho Er SYSTEMS

13

of the alloys are noticeably larger (see Table IV). Although the magnetic properties of Dy and a 50 at.% Ho in Tb alloy have been observed to be quite similar24. and the thermal expansion versus temperature curves of Dy and a 50 at.?, Ho in Tb alloy have been shown recently to be very similar 25. the deviation in c does indicate some difference between Dy and a 50 at.“, Ho in Tb alloy. ACKNOWLEDGMENTS

The authors would like to thank P. E. Palmer and J. J. Croat for making the pure metals, J. Moorman for running some of the DTA experiments. H. Hensch for assistance with the X-ray studies and H. Baker for the metallographic examination of the alloys. The careful mass spectrometric analysis of the metals by R. Conzemius and vacuum fusion analysis by N. Beymer are also greatly appreciated.

REFERENCES I 2 3 4 5 6 7 X 9 10 II 12 I3 I4 I5 I6 I7 IX I9 20 21 22 23 24 25

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