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AD-760 777

MATERIALS PROCESSING OF RARE EARTH COBALT PERMANENT MAGNETS

Paul '. Jorgensen, et al

Stanford Research Institute

Prepared for:

Air Force Materials Laboratory

February 1973

DISTRIBUTED BY:

KFÜ National Technical Information Service U. S. DEPARTMENT OF COMMERCE 5285 Port Royal Road, Springfield Va. 22151

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AFML-TR-73-60

MATERIALS PROCESSING OF RARE EARTH COBALT PERMANENT MAGNETS

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P. J. Jorgensen R. W. Bartlett

Stanford Research Institute

TECHNICAL REPORT AFML-TR-73-60

February 1973

Approved for public release; distribution unlimited.

Reproducpcl by

NATIONAL TECHNICAL INFORMATION SERVICE

U S Daparlmtnl of Comirnr« Si)fing(i»l'l VA 771 "it

Air Force Materials Laboratory Air Force Systems Command

Wright-Patterson Air Force Base, Ohio 45433

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NOTICE

When Government drawings, specilications. or other data arc used lor any purpose other than in connection with a definitely related Government procurement operation, the United States Government thereby incurs no responsibility nor any obligation whatsoever; and the tact that the govern- ment may have lormulated, lurnished, or in any way supplied the said drawings, specilications, or other data, is not to be regarded b" implication or otherwise as in any manner licensing the holder or any other person or corporation, or conveying any rights or permission to manufacture, use. or sell any patented invention that may in any way be related thereto.

Copies of this report should not be returned unless return is required by security considerations, contractual obligations, or notice on a specific document.

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üNCa^ASSIFIED

MATERIALS PROCESSING OK KAItE EAHTH-COHALT PERMASEMT MAGNETS

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Paul J, Jorgenaen Robert w. Bartlett

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February 1973 •• CONTMACT OR CtANT NO

F33615-7()-C-162 1 h PHOJ I C T NO

ARPA Order No, 1617

Program Code No, OD10

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Approved for public release; distribution unlimited,

.PPLLME». '*Ni . IP ON SOn IN 6 Mil 'A^*^ ACTIVITY

Air Force Materials Laboratory Wrißht Patterson Air Force Base Ohio 15433

* ÜST«, A C

A study of the oxidation kinetics of SmCo involving selective oxidation of

samarium was extended from high temperatures to lüo'C without a change in the diffud.01

controlled rate mechanism. The low temperature oxidation rates are consistent with

the amount of oxidation normally observed in SmCo powders and with the temperature

rise expected from oxidation of ground powders. The solubility of oxygen in SmCo5 at

1130OC (sintering temperature) in excess of the 800 C solubility was determined to be

3500 + 500 ppm. The source of oxygen during sintering is the oxide subscale. Sub-

micron oxide particles precipitate within the SmCor grains on cooling from the sinter-

ing temperature. Oxidation causes depletion of samarium and precipitation of Sm^Co

particles within grains. Both inclusions are postulated to be sources of domain wall

nucleatlon. The oxide inclusions can be removed at 800 C by an aging treatment, that

collects the oxide into a few large grains ouUlde the SmCo5 grains by a solution-

reprecipitation mechanism involving grain bounoary transport of samarium. The Sm2Co

17

inclusions are not affected by thermal aging, but they can be prevented from occurring

by including excess samarium in the sintered compact to replenish samarium lost by

oxidation.

The kinetics of solid phase sintering of SmCo have been studied as a function of The kinetics can be described by a grain time, temperature, and stoichlometry

md oy

boundary transport mechanism that Is independent allr"

the samarium cobalt ratio In llie

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Magnetic Particles

Rare Earth

Cobalt

Oxidation

Internal-Oxidation

Sintering

Diffusion

Magnetic-coercivity

DD .^..1473 'BACK, (PAGE 2) /£ UNCLASSIFIED

Security Classification

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MATERIALS PROCESSING OF RARE EARTH COBALT PERMANENT MAGNETS

P. J. Jorgensen R. W. BartlfU

Approved for public release; distribution unlimited.

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FOREWORD

This is the fifth semiannual interim technical report of the research program "Materials Processing of Rare Earth-Cobalt Permanent Magnets" under Contract F33615-70-C-162A. Stanford Research Institute project number is PYU-8731. This project is being conducted by the Materials Laboratory of Stanford Research Institute. Dr. Paul J. Jorgensen, Manager of the Ceramics Group, is the project supervisor. Dr. Robert W. Bartlett of Stanford University is project consultant. The research described in this report is part of the contractual research program of the Electromagnetic Materials Division, Air Force Materials Laboratory, Air Force Systems Command, Wright-Patterson AFB, Ohio. Mr. Harold J. Garrett (AFML/LPE) is the project monitor. It was sponsored by the Advanced Research Project Agencv, ARPA Order No. 1617, Program Code No. OD10.

This report covers research conducted between July 1 and December 31, 1972 and was submitted in February 1973 by the authors for publication.

This technical report has been reviewed and is approved.

/^/^«^v^^V

CHARLES H. ROBISON, Major, USAF Chief, Solid State Materials Branch Electromagnetic Materials Division Air Force Materials Laboratory

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ABSTRACT

A study of the oxidation kinetics of SmCo Involving selective oxi- 5

dation of samarium was extended from high temperatures to 100 C without

a change in he diffusion-controlled rate mechanism. The low temperature

oxidation rates are consistent with the amount of oxidation normally

observed in SmCo powders, and with the temperature rise expected from

oxidation of ground powders. The solubility of oxygen in SmCo at 1130 C 5

(sintering temperature) in excess of the 800 C solubility was determined

to be 3500 ± 500 ppm. The source of oxygen during sintering is the oxide

subscale. Submicron oxide particles precipitate within the SmCo grains on 5

cooling from the sintering temperature. Oxidation causes depletion of

samarium and precipitation of Sm Co particles within grains. Both

inclusions are postulated to be sources of domain wall nucleation. The

oxide inclusions can be removed at 800 C by an aging treatment, that

collects the oxide into a few large grains outside the SmCo grains by a D

solution/reprecipitation mechanism involving grain boundary transport of

samarium. The Sm Co inclusions are r.ot. affected by thermal aging, but

they can be prevented from occurring by including excess samarium in the

sintered compact to replenish samarium lost by oxidation.

The kinetics of solid phase sintering of SmCo have been studied as 5

a function of time, temperature, and stoichiometry, The kinetics can be

described by a grain boundar' transport mechanism that is independent of

the samarium/cobalt ratio in the alloy.

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CONTENTS

I INTRODUCTION 1

II OXIDATION 0"' SmCc) 3 5

A. Oxidation, Subscale Structure, and Composition 3

B. Kinetics of Internal Oxidation of SmCo,. at Elevated 5

Temperatures 5

C. Kinetics of Oxidation at Low Temperatures . 8

D. Expected Temperature Rise Associated with Low Temperature

Oxidation 14

E. Oxidation Rate Mechanism 16

III VACUUM ANNEALING PARTIALLY OXIDIZED SmCo AT SINTERING

TEMPERATURES 25

IV OXYGEN SOLUBILITY DURING SINTERING AND SUBSEQUENT OXIDE

PRECIPITATION IN SmCo^ 31 5

V PRECIPITATION OF Sm Co IN SmCo 39

VI DISCUSSION OF Sm 0 AND Sm Co INCLUSIONS AND INTRINSIC

COERCIVITY IN SINTERED MAGNETS 43

A. Summary of Inclusions 43

B. Inclusions and Intrinsic Coercivity 43

VII SINTERING

A. Introduction 49

B. Experimental Procedure 51

C. Results and Discussion 52

REFERENCES . 63

Preceding page blank

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ILLUSTIlfVriüNS

1 Course Fibrous Structure of Samarium Oxide in the Oxide/

Cobalt Composite Subscale Grown at High Temperatures, 1150OC 4

2 X-Ray Diffraction Spectra lor the Oxide Subscales Grown in

Air and in Pure Oxygen 6

3 Arrhenius Plot for the Internal Oxidation Rate at High Temperatures 7

1 OxyRen Dissociation Pressures for CoO in a 90% CO + 10% CO 2

Mixture 9

5 Internal Oxidation Rate Dependence on Oxygen Pressure (Air

Equilibrium at Surface) 10

6 Internal Oxidation Rate Dependence on Oxygen Pressure

(CoO Equilibrium at Surface) 11

7 Thermogravimctric Data for Internal Oxidation of 25 ^m SmCo Particles 13

8 Arrhenius Plot for the Internal Oxidation Rate Extended to Lower Temperatures 15

9 Ad_abatic Temperature Rise Accompanying Partial Oxidation to Sm 0 17

6 10 Oxygen Self-Diffusion in B-Type Sm 0 (After Stone ) .... 19

11 Fibrous Samarium Oxide with a Crack in the Subscale Grown at 800OC 22

12 Oxide Depletion Zone Resulting from Internal Oxidation of

SmCo Followed by Vacuum Annealing, 125X 26

13 Oxide Depletion Zone at Higher Magnification (SEM) 27

14 Boundary of the Oxide Depleted Zone (Right Side) Within the Subscale 28

vi

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ILLUSTllATIONS (Concluded)

18 Boundary Between the Oxide Depleted Zone and the SmCo Core 30 o

16 Oxide Precipitates in the SmCop Grains after OxyRen Satura- tion at 11250C and Asin« at 8o80C 33

17 Quenched and Aged SmCo Grains 35 D

18 Samarium Oxide Particles in a Sintered SmCo,. Magnet Uncovered 5

by Etching 37

19 Scanning Electronmicrograph of Sm Co Precipitates in SmCo 40 ^ 17 O

20 Magnetic Domains Passing Through Sm Co Precipitates Within

SmCo. Grains 16 5

21 Linear Thermal Expansion of Polycrystalline Randomly

Oriented SmCo^ 53 5

22 Samarium-Cobalt Shrinkage Isotherms for an Alloy Contain ng

27.7 wt% Sm 51

23 Samarium-Cobalt Shrinkage Isotherms for an Alloy Containing

31.35 wt% Sm 55

24 Samarium-Cobalt Shrinkage Isothems for Alloys Containing

33.0 and 33.33 wt% Sm 56

25 Arrhenius Plot of the Shrinkage Rate Versus Reciprocal

Temperature 58

TABLES

1 Some Efects Involving Magnetic Coercivity in SmCo5 and

Their Possible Causes 'I7

2 Sintering Constants for Equation 4 50

3 Oxygen Analysis of SmCo Alloys 59

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I INTRODUCTION

The objectives of this program are to imestigate materials pro-

cessiriK methods with the goal of achieving optimum properties of SmCo 5

and other RECo compounds as permanent magnets. The primary problem

is achieving high magnetic coercivity.

Previous to this research period, .1 variety of unconventional pro-

cesses for generating RECoc powders were investigated but none was found

to be more effective in obtaining high coercivity materials than the

properly conducted comminution process of bull milling.

The sintering process is also a critical step in magnet production.

Studios to gain further knowledge of the liquid-phase sintering process

and magnetic evaluation of sintered alloys were carried out prior to this

research period, and studies of ehe final stage of sintering (solid-state

process) were conducted during this period and arr reported.

Selective internal oxidation of samarium invariably accompanies

magnet powder processing. A study of the kinetics of oxidation and the

morphology of the oxide subscale was completed during this period. Studies

on oxygen solubility in SmCor and the precipitation of oxide inclusions D

during quenching and aging are reported.

A study of Sm Co precipitation in samarium-depleted SmCo grains is j i # 5

included in this report.

—■-"- ■-- —

II OXIDATION OF SmCo 5

A. Oxidation, Subscule Structure, and Cumpositlon

Oxidation of SmCo5 is selective, resulting in a subscale consisting

of samarium oxide platelets ana libers surrounded by cobalt. The oxide

fibers are oriented in their growth direction, which is generally normal

to the alloy surface. Scanning electron micrographs on the subscale

structure in the direction of growth and transverse to the direction of

growth are shown in Figure 1. These specimens were oxidized at USt/'c

lor two minutes in a 90% C02 + 10% CO mixture, which prevents any oxidation

of cobalt. After motallographic mounting, the specimens were etched in

50% HN0:} + 30% H2S04 + 20% ^0 and vapor coated with a gold-platinum alloy.

The etching removes cobalt and leaves the oxide structure in relief.

The transverse section indicates that the oxide consists of fibers, narrow

platelets, and a loosely Interconnected network of joined platelets.

Previous x-ray diffraction determinations regarding the crystallo-

graphic structure of the composite subscale formed by internal oxidation

of SmCo5 have been somewhat confusing. More detailed x-ray diffraction

studies following oxidation have now been completed on both powder samples

and flat surfaces )f SmCo inpeots. These studies have been made in pure

oxygon, in air, and in the 90% C02 + 10% CO mixture. Determinations have

been nade on samples oxidized at temperatures from 320 to lOOo'V In all

casej the primary metal produced is fl-cobalt (cubic).

The samarium oxide phase that is obtained during oxidation is either

Sm203 0r an oxynitri(lG depending on the presence or absence of nitrogen.

B-type Sm^^ (monoclinic) is produced at all oxidation temperatures. Pure

Sm 0 exhibits the monoclinic structure above a50-900OC, and C-type Sm 0 * 0 2 3

Preceding page blank

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(a) TRANSVERSE

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FIGURE 1 COARSE FIBROUS STRUCTURE OF SAMARIUM OXIDE IN THE OXIDE/COBALT COMPOSITE SUBSCALE GROWN AT HIGH TEMPERATURES,

1150'C

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(cubic) is usually observed at lower temperatures. X-ray diffraction

peak intensities lor the C-type sosquioxide are very weak or nonexistent.

No other samarium oxide phases are observed in the absence ol nitrogen.

After oxidation in air, a suboxynitrido phase, SmN 0 (NaCl structure) 1-x x '

is .ibservod and the Sm 0 peaks are weak. This samarium compound lias been 9

previously prepared and idcntilied by Kelmlee and EyriiiK." Electron-beam

microprobe (EDM) analyses on air-oxidized samples have indicated that the

actual oxygen content oi the subscale is somewhat less than that required

stoichiomotrically for Sm 0 and somewhat greater than that required

stoichiometrically for SmO, This finding is consistent with the presence

of both the sesquioxide and suboxynitrido. However, nitrogen was not

detected by EBM analysis.

X-ray diffraction spectra taken after oxidation of subscale at 1000OC

in air and oxygen are shown in Figure 2. The shift from the sesquioxide to

the suboxynitride caused by nitrogen in the air oxidation tests is clearly

evident.

B- Kinetics of Internal Oxidation of SmCo_ at Elevated Temperatures ' O ' .....

The oxide subscale thickness increases parabolically with time

indicating a diffusion-controlled rate-process. The kinetics of this

process for both SmCor and PrOo are identical and have been studied at

elevated temperatures using a metallographic technique on small sectioned

cast pieces.

The temperature dependence of the internal oxidation rate follows

the Arrhenius law. Further work at higher temperatures conducted during

this period has shown that there Is a break in the Arrhenius plot, with

a lower apparent activation energy for the oxidation rate at higher

temperatures. The break in the curve occurs at about 65ü(>C (see Figure 3).

The apparent activation energies are 14.0 kcal/mole at lower temperatures

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and 7.7 kcal/mole at higher temperatures. Each dat"ni point shown in

Figure 3 represents a .riplicate set of runs.

With two exceptions, all the data shown in Figure L are from runs

conducted in one atmosphere of air or oxygen. The remaining two data

points were obtained from runs conducted in a flowing gas mixture of

mr CO + 101 CO. The chemical potential of oxygen in this gas mixture 2

is below that required for oxidation of cobalt; see Figure 4. Conse-

quently, the usual thin oxide scale containing cobalt oxide over the

subscalc is not formed.

The markedly reduced oxygen potential in the 90% C02 + 10% CO mixture o

caused a small decrease in the subscale oxidation kinetics at 840 C and

nt J25"c. In both cases the rates wen; reduced to about half the respective

rut.'s obtained at the same temperatures In air. As determined from these

limited data, the parabolic oxidation rate dependence on oxygen partial

pressure is summarized in Figures 5 and 6. If the exterior samarium oxide

in the subscale is in equilibrium with air at one atmosphere, then the

increase in the parabolic rate constant with oxygen pressure is shown by

Figure 5. However, if the exterior samarium oxide in the subscale formed

by air oxidation is assumed to be in equilibrium with CoO, the oxygen

pressure dependence is shown by Figure 6. In both cases the oxygen pressure

dependence is very weak. Generally for parabolic oxidation of metals, it

is related to the point-defect transport mechanism in the oxide. Although

not enough is known about the defect mechanism in rare earth sesquioxides

to draw firm conclusions, the observed oxygen pressure dependence is less

than that encountered for any of the usual defect mechanisms responsible

for ion diffusion in sesquioxides.

C. Kinetics of Oxidation at Low Temperatures

Oidinarily, the oxidation of magnet powders occurs during handling

at room temperature, or at slightly elevated temperatures resulting from

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TEMPERATURE

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FIGURE 4 OXYGEN DISSOCIATION PRESSURES FOR CoO IN A 90% C02 + 10% CO MIXTURE

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FIGURE 5 INTERNAL OXIDATION RATE DEPENDENCE ON OXYGEN PRESSURE (AIR EOUILIBRIUM AT SURFACE)

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FIGURE 6 INTERNAL OXIDATION RATE DEPENDENCE ON OXYGEN PRESSURE (CoO EQUILIBRIUM AT SURFACE)

1 1

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the exotliermic oxidation reaction. Therefore, it is particularly

important that kinetic data bo obtained at very low temperatures so

that the extent of oxidation as a function of particle size, time, and

temperature can be predicted. Previous attempts to measure small amounts

of oxidation using a zlrconiu electrochemical cell, which is very sensi-

tive to small amounts of oxygen consumption at low oxygen pressures, 3

have been reported but the results were not very satisfactory. A review

of those data after accumulating more knowledge about ^idation at higher

temperature indicates that the ma.jori'y of these runs represent measure-

ments in such small amounts of oxygen that adsorption and desorption

effects were probably dominant. As a consequence, additional oxidation

kinetic studies with the zlrconia coll apparatus have been conducted in

higher oxygen pressures, although the oxygen pressure is still sufficiently

low that differences in oxygen pressure can be measured and used to o

detemine the oxidation kinetics below 200 C.

An electronic microbalance has been used to measure the oxl'' .ion o

kinetics gruvimetrically in the temperature range from 150 to 250 C. The

tliei-mogravimetrie data are shown in Figure 7.

For both the gravimetric und oxygon consumption methods, the results

have boon obtained using a single batch of SmCor particles with an average 0

particle diameter of 25 micrometers. This diameter was determined by

photographing and measuring numerous particles and then averaging tho diame-

ter. In those experiments the extent of oxidation was limited to an oxide

scale thickness very much smaller than tho particle radius. Conseqoontly,

one dimensional planar diffusion was an adequate approximation in calcu-

lating tho diffusion-controlled oxidation rate.

Tho weight gain and oxygon consumption curves obtained from these

experiments wore parabolic. The results were converted into equivalent

oxide scale thickness rates for comparison with tho oxide scale thickness

rates determined by the sectioning technique used at higher temperatures,

12

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FIGURE 7 THERMOGRAVIMETRIC DATA FOR INTERNAL OXIDATION OF 25-//m SmCot PARTICLES

13

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The oarabolic oxidation rates for the low temperature experiments

are shown in Figure 8, The straight line is an extrapolation of the

Arrhenius plot from the high temperature data of Figure [i. These data

show that the same oxidation mechanism observed at higher temperatures

continues down to the lowest temperature at which rates can be measured,

100 C. The kinetics for the oxidation process shown in both Figures 3

and 8 span a range of oxidation rates that is seven orders of magnitude,

o o between 100 C and 1125 C.

U, Expected Temperature Rise Associated with Low Temperature

Oxidation

The adiabatic temperature rise caused by partial oxidation of SmCo 0

has been calculated. However, the assumptions for this calculation need

to be emphasized. First, samarium is selectively oxidized to samarium

sesquioxide and hence all of the measured oxygen in a sample can be

assigned to Sm 0 . Second, the sensible heat of oxidation is uniformly

distributed in the alloy particle instantaneously; rapid heat transfer

occurs in the metal and the particle surface temperature, wheve oxidation

takes place, is no greater than the interior temperature. This approxi-

mation is probably adequate for a moderately slow oxidation process and

small particles such as those Involved in preparing sintered magnets. The

third assumption is that no heat is lost from the oxidizing metal system;

the aggregate alloy powder is an adiabatic system. This approximation is

probably valid for a large volume of particles and oxidation occurring

fast enough that heat transfer out of the crucible or particle container

is not significant.

The temperature rise will be proportional to the amount of oxygen,

the specific heat of SmCo , and the heat of formation of Sm O , The 5-l -1 2 J

specific heat is 0.1 cal g deg , which was estimated from an assumption

-1 -1 that the heat capacity is 7 cal g-atom deg . The heat of formation of

o 4 Sm O, at 298 K is 433,900 cal/molc,

14

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15

Experimental:

• ■ Microbalance

A - Oxygen Consumption —

FiGURE 8

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ARRHENIUS PLOT FOR THE INTERNAL OXIDATION RATE EXTENDED TO LOWER TEMPERATURES

"-'••'-' '■;

I III lIHPILWWV^nV* - ",1" ll"1 ' '■ "' l-. — .-F»-™.^.-. ......... ... „.»«». ..... —r. mi^^tPWip^lH l||l| I ■■ l»l»»MW»»tT»'«l». I I I IIWII

The temperature rise caused by an increase in oxygen is shown in

Figure 9. An increase of 1000 ppm of oxygen causes an aciiabatic tomper-

o ature increase of nearly 100 C.

5 As previously discussed, oxygen analyses in samarium cobalt materials

after grinding generally range from a few hundred to a few thousand ppm

depending on the processing and particle size. This amount of oxygen is

sufficient to account for an appreciable adiabatic temperature rise

o o during oxidation. A 100 C rise in temperatux-e from 25 to 125 C is suf-

ficient to account for a subscale thickness of about 100 A in a few hours.

o Furthermore, a 100 A subscale thickness is equivalent to a 1000-ppm oxygen

* increase for a 3-micrometer-diameter SmCo,. particle. Hence, for the

5

particle sizes employed in sintered magnets, the observed increase in

oxygen is consistent with the particle size, the expected temperature rise,

the kinetics of oxidation, and the expected subscale thickness as a result

of the temperature rise. Although these relationships do not prove that

a temperature rise and the internal oxidation by the diffusion-controlled

mechanism have occurred to match the observed oxygen concentration, they

at least show that all of the observed results are compatible.

E, Oxidation Rate Mechanism

3 It was previously shown that the published data on the oxygen

solubility in cobalt coupled with the rapid rate of internal oxidation

required that the diffusion coefficient for oxygen In cobalt would have

-6 2 o to be 1 x 10 cm /sec at 420 C and even greater at higher temperatures.

In fact, at high temperatures the diffusion coefficient must be greater

-4 -6 2 than those typically encountered for liquid metals (10 to 10 cm /sec).

Consequently, it appeared that the transport of oxygen in the B-cobalt

matrix could not adequately explain the rapid rate of growth of the

subscale.

* See Figures 27 and 28 of Reference 3,

16

■—■ UMMi^ ^^mmrn^. ^^^^M^^^.

^™r"-""'FrTF<™w«»w™«»'!*'™''«'™"'!"»r™w<^

200

0 200 400 600 800 1000 1200 1400 1600 1800

OXYGEN — ppm ,. „„, nn rA-8731-e2

FIGURE 9 ADIABATIC TEMPERATURE RISE ACCOMPANYING PARTIAL OXIDATION TO Sm203

17

mmm^^M,, . tmf'---~~ - ■

^JfOIMMI in i iipiMia '■ >'»iii'i mi ■ IM--Ii' . i- IIIWIMILimil ivi ' ^^ W11P»WI«111I 11 ■■•■

A literature search to determine the available information on

diffusion of oxygen in B-type and C-type rare earth sesquioxides has

been made and it presently appears that oxygen anion transport within

the oxide fibers is also too slow to account for the rapid growth of

the subscale. Kofstad provides an adequate review of the existing

diffusion literature for the rare earth oxides.

A study of self-diffusion of oxygen in D-type Sm 0 has been made

6 o by Stone over a temperature range between 693 and 907 C using an isotopic

exchange experimental technique. The activation energy for diffusion

was about 23 kcal/mole, and the data are summarized in Figure 10. Wirkus,

7 8 Berard, and Wilder ' have rcoxidized partially reduced rare earth

sesquioxides, and although they have not studied Sm 0 , they have obtained

diffusion coefficients of the same order of magnitude in the same tempera-

ture ranges as the diffusion coefficients determined by Stone. Their

implicit assumption is that these oxides are semiconductors, whereas Tare

9 and Schmalzried find them to be ionic conductors at low oxygen pressures

o (C-type structure). At 800 C the diffusion coefficient of oxygen in

-10 2 -1 Eu 0 determined by Wirkus and Wilder is D = 2.27 x 10 cm sec .

2 3 0 The value obtained from Stone's data for oxygen diffusion in Sm 0 at

o -10 2 -1 800 C is D = 2.76 x 10 cm sec .

0

If the rate of internal oxidation of SmCor is determined by the flux 5

of oxygen anions through Sm 0, platelets, then one can deduce the product

D AC from the parabolic rate constant. It is instructive to make this

calculation on the assumption that oxygen transport through the oxide is

rate-controlling at 800 C, and compare the results with the diffusion data

at the same temperature. The derived value of D AC using this hypothesis -7 -1

is 1.67 x 10 g cm sec . Matching the value for the diffusion coef- -10 2 -1

ficient determined by Stone, namely D = 2.76 x 10 cm sec , would 3

require an oxygen concentration difference of about 1000 g/cm . This

conclusion is clearly absurd since the total oxygen concentration is only

18

MHMMM I^IIM itmiiiiiMi v i - -■--^-'liiiiir'rfilliii'üllinifi'ill'iil i

• ii w»"r^p^^p^B»pi^w^'i''■" ' ' l,lll" m*mi*\vmmn^mm^w*~>*'i^^mmimmmi*r***'si'*'^ii**imimm***mmimi^t^mi^mi^mmm'^^^

10 8

10'

io-10u-

E

Q0

10 -11

10 12

= 1 1 1 1 1 1 z

1\ ■ ■

I \\

- \ \ \\

= >v\

—

' \

k 99 999% Sm203 -

\ — 99.9% Sm203 I -

._ \

....

— W —

— —

w \\

— \\ — \\ — \\ — w — — -

\

1 1 1 1 1 1 0.6 0.8 1.0 1.2 1.4

103/T — V1

1,6 1.8 2.0

TA-8731-81

FIGURE 10 OXYGEN SELF-DIFFUSION IN B-TYPE Sm^ (AFTER STONE6)

19

JiMtiiiiiiiiHiMmiiiBii II MwyiMM

"*'■> ' ■ ■■ ■ »i 11 ii i in mmmmmmfmwmmm^i^r^*^^****'*n ' "vmmmmmrmnrmr'mm i i mmmtw^mmmmirmiimfrtm

1.15 g cm and thu actual driving force could only represent a concentra-

tion difference that is matched by an appropriate point-dofect concentra-

tion difference In the oxide that would be undoubtedly much smaller than

the total oxygen concentration in the oxide. It is quite clear from these

calculations that if oxygen transport through the oxide is the rate-limiting

step, then the diffusivity must be much faster than the diffusivity observed

by Stone and by other investigators involved In the diffusion of oxygen

or reoxidation of partially reduced B-type and C-type rare earth sesqui-

uxides.

Vacuum annealing a partially oxidized sample at a temperature higher

than the oxidation temperature evidently prevents further oxidation when

the sample recnters oxygen. As discussed in the next section, annealing

causes dissolution of tue aligned fibers in the subscale and reprecipitation

of larger, more equidimensional oxide particles. The samarium oxide par-

ticles remain as the dispersed phase in the subscale while 0-cobalt and,

additionally. Sm Co are the matrix phases as a result of high temperature

vacuum appealing.

If rapid oxygen transport through fl-cobalt were rate-controlling,

then continued oxidation would be expected after vacuum annealing and

reexposing the specimen to oxygen at elevated temperatures. Hence, oxygen

diffusion in the cobalt phase is eliminated as the rate-controlling step.

The dispersed oxide phase occupies approximately 40 percent of the

subscale volume and there is scanning microscopic evidence that considei'-

able linking (sintering) exists among the various oxide particles after

vacuum annealing the subscale. Hence, if oxygen anion transport through

the oxide were rate-controlling, it is expected that there would be some

continued oxidation rather than passivation.

There are no consistent phase changes in either cobalt or the

samarium oxides that correlate with the apparent change in activation

20

— - - -*-^-- ■ -■-■-"--"-■■'—

r~"—'— i. i nm^^^mm m unjii ii mmmmmpimmmsn***si. •w.'^m^immmm^i^w^nKmimwfmmmiK^^'maßfm^m^^f^rmmmrmrmmi^mmmmmmimfm

energy above Toi/'c. This supports the conclusion that bulk diffusion of

oxygen in neither cobalt nor Sm 0 is fast enough to account for the

internal oxidation rate.

Migration of samarium outward by a vacancy exchange mechanism would

cause the accumulation of internal vacancies and the precipitation of voids,

which are not observed.

The rapid oxygen transport required for internal oxidation may occur

along the interface between oxide fibers and ß-cobalt. Because of the

small dimensions of the composite microstructure of the subscale, a large

amount of interface area is present. This mechanism can account for the

apparent change in activation energy for the internal oxidation rate

at about 700OC. The fiber size and spacing begin to increase with

increasing temperature at about that temperature and the interfacial area

per unit volume of oxide subscale decreases correspondingly. This effect

is shown by comparing Figure 1 (1150OC) with the scanning electron micrograph

shown in Figure 11 for subscale oxidation at 800 C, If it is assumed that

the same transport mechanism is involved at all of the temperatures for

which there are experimental data, then a reduction in the interfacial area

would reduce both the number of effective conduits for oxygen transport and

the apparent activation energy, even though the true activation energy for

surface diffusion and the enthalpy for generating mobile species at the

samarium oxide-cobalt interface remain constant.

Other experimental techniques for elucidating the mechanism of oxygen

transport involved in internal oxidation of SmCo have been attempted

without much success. Ion microprobe analysis following oxidation of SmCo5

18 16

with 99% 0 was not successful because considerable amounts of 0 wore

also observed. The source of the O is not clearly known, but evidently

a considerable amount of oxygen and water vapor adsorb on the sample and

is subsequently desorbed during ion microprobe analysis. This contaminating

21

fa ,.■„,■,...: ,.....„ i II Vi'itilttlrtMntii i -- ------ '

- "' ■■'■■■■■ "-p"»i—■■ ■• mim mit^mn^^i in i •mm*mmm*mmmmmm I I I I IIHU IMMM^p

TA-8731-91

FIGURE 11 FIBROUS SAMARIUM OXIDE WITH CRACK IN THE SUBSCALE GROWN AT SOCfC

22

-"■>—*-**^-*— ■Mil I1M IHMTiMn mi

.« I. II|I.HI^ I fc,*"*-PT^ .+,.:.»«..—H' i -r— ' -—I^II m u ..1U1|1I||IUII1II- _ , ^PWH^^»«^»*^^'»-™»»"!"^ ^^»

oxygen is probably adsorbed during the transfer of the sample from the

0 furnace chamber to the ion microprobe. Adsorbed oxygen and water

vapor may continuously migrate to the ion impact area by surface diffusion

and then be discharged and picked up by the ion detector. Although

18 16 moderately high levels of 0 could also be detected with the 0 con-

tamination, it was not possible to do marker experiments unambigously,

18 18

using 0 and subsequently detect the location of the 0 in the subscale.

Auger spectroscopy was also employed to investigate the oxide films

formed by low temperature oxidation. These experiments also indicate that

a considerable amount of adsorbed species, some of which contain oxygen,

are migrating to the surface and affecting the results. In both ion probe

analyses and Auger analysis, the oxygen concentration pivfile obtained from

continued sputtering appears to extend deeper into the specimen than is

known to be true from kinetic and metalloeraphic data.

23

•~~rr~—~'~~~—~—~~^~—-r^~*mm^^*m^~'—~~~'^^~^~^m^w-m~^^m^mm^mmm^m

III VACUUM ANNEALING PAKTIALLY OXIDIZED SmCo AT SINTERING TBiraRATURES

An example of the result of high temperature annealing in the

absence of oxygen after partial internal oxidation of SmCo is shown 5

in Figure 12. This particular specimen was oxidized 15 minutes in air

o . o at 800 C, and was subsequently annealed 16 hours at 1125 C. A scanning

electron micrograph of the same section is shown in Figure 13. The

original subscale prior to vacuum annealing penetrated to the farthest

extent of the large oxide particles shown in Figure 12 and had the

fibrous microstructure shown in Figure 11. As a result of vacuum anneal-

ing, the fine samarium oxide fibers have recrystallized into large,

roughly equidimensional particles. The subscale contains a conjugate

zone partially depicted in samarium oxide adjacent to the original SmCo 5

core. The few oxide particles in the oxide depleted zone are consider-

ably larger than the oxide particles in the undepleted part of the

subscale. A higher magnification view of the reprecipitated oxide

particles Is shown in Figure 14.

The oxide particles and alloy phases have been identified by electron

beam microprobe analyses. The matrix of the depleted region consists of

Sm D) . This phase results because samarium is transported from the 2 17

SmCo core into the subscale region where it reacts with cobalt. The 5

SmCo core is transformed from a single phase region into an SmCo matrix 5 5

containing Sm Co precipitates as a result of samarium depletion from

the core. The solid state diffusion and precipitation reactions are:

25

Preceding page blank

■mi

-OXIIH ■

$ - Co-

SmCor > 5;

AFTER OXIDATION

ii - Co -V* SmjCo^ —

OXIDE ■

OXIDE DEPLETED ZONE

SmCo

(Sm 5 ) AFTER

2C())7i)t)il J ANNEALING

&' I >

TA 8731 85R

FIGURE 12 OXIDE DEPLETION ZONE RESULTING FROM INTERNAL OXIDATION OF SmCo5 FOLLOWED BY VACUUM

ANNEALING, 125X

LT.

_-,

FIGURE 13 OXIDE DEPLETION ZONE AT HIGHER MAGNIFICATION (SEM)

27

SA 8731 93

FIGURE 14 BOUNDARY OF THE OXIDE DEPLETED ZONE (RIGHT SIDE) WITHIN THE SUBSCALE

2*

(1) At vacuum annealing temperature, diffusion reaction:

core + subscale matrix - subscale matrix + core

Mfy + ^ ^ - f ■■a0oi7 + ^c-x^s <1)

where 6 ~ 1.

(2) During quenching, precipitation:

If 6 = 1, then

Smi-xC05 - ( ' " T1 ) SmC05 + ( ? X) Sm2C017 ' (2)

While samarium migrates out of the core, oxygen dissolves and

mlRrates Into the core, causing the reduction In population of oxide

particles In the depleted zone. This phenomenon Is a direct result of

the solubility of oxygen In SmCo at the higher vacuum annealing temper- o

o aturc, 1125 C, Although the vacuum annealing time In this case was

16 hours, Identical results have been obtained with a 3Ü-mlnute vacuum

anneal, which Is Identical to our standard magnet sintering time and

temperature.

The scanning electron micrograph shown in Figure 15 shows essentially

all of the features of this system. The original boundary separating the

subscale from the SmCor core Is shown. Uelow this boundary Is the region

containing a Sm Co matrix and large samarium oxide parllcles. A two-phase

alloy region consisting of a SmCo,. matrix and Sm Co precipitates exists 5 2 17

above the boundary. A relic of the original fibrous subscale is also shown

in this unusuaJ SEM micrograph. A small samarium oxide particle is also

shown in the core region above the boundary line. This particle is pro-

duced by oxygen dissolved in SmCo and subsequently precipitated on cooling. 0

29

____

IT i .11 IPf

SA 873! 94

FIGURE 15 BOUNDARY BETWEEN THE OXIDE DEPLETED ZONE AND THE SmCob

com

30

imaum ■ M^——■ I .

■ipiiyw^piiniliiuiini iiii.iiii .mimUM-tiLUWIW. ^.»PI'K^'^'^WBWWiPiPWWffWWI'll'lll WWliH'UWUWl.f ■II.'II.'WIWIPI mm i .•■.•••■•.•« li nai i 11 ii ii inn i mi ill II I ,I«|IBIP|

IV OXYGEN SOLUBILITY DURING SINTERING

AND SUBSEQUENT OXIDE PRE-

CIPITATION IN SmCc) 5

Attention is now turned to the problem of oxide dissolution in

SmCo at sintering temperatures and the subsequent redistribution of the 5

oxide on cooling.

The experimental procedure used in this study has been to introduce

considerable amounts of oxygen deliberately into the SmCo system using o

small polycrystalline cast pieces by first partially oxidizing the

relatively oxygen-free alloy. Because of their large grain size and

absence of porosity, the cast alloys are easier to examine metallograph-

ically by SEM and by EBM methods than sintered magnets. An oxidation

temperature of 800 C was chosen because it is substantially below the

sintering temperature, and the homogeneity range in SmCo is believed 5

to be extremely narrow at that temperature. Furthermore, it has been

shown that aging sintered SmCo. magnets at about this temperature (usually 5

900 C) can often be beneficial to the magnetic properties. In particular,

10 an increase in magnetic coercivity is often obtained. After oxidation,

o the specimens were vacuum annealed at 1120 to 1125 C for 30 minutes.

We were initially concerned about obtaining chemical equilibrium through

the entire ingot, which is typically 0.5 cm on edge, and therefore very

long annealing times were used. Subsequently, we found that the standard

sintering time, 30 minutes, is sufficient to provide a macroscopically

homogeneous distribution of oxygen in the SmCo core of such large alloy 5

pieces. The initial alloy had the stoichiometric SmCor composition, i.e., D

there was essentially no depletion of samarium.

The results will be described for specimens after vacuum annealing

at sintering conditions and also after a third heat treatment, i.e,

31

■'■■ :■•■■■ --.■■-i.-..-;n. ,■

uttm ihlini ■"-—-'"-•J-'; --- - ■ "'— --

|lliHjmM>umiiiiuw.iwii.|ii,iuiiw^ HI «iMjimnuwpjiiiM << ' ■«■■^^■^ > m« ^*mmmm

aging at 80Ü C for 45 minutes. For convenience wo will refer to the

two types of specimens as "quenched" and "aged," Quenched specimens

are produced by quickly removing the specimen from the furnace, but

cooling occurs in vacuo so that "normalized" would be a more appropriate

classical metallurgical term.

The amount 01 oxygen transferred from the oxide subscale into the

SmCo was crudely estimated by measuring the thickness and the area of 5

the depletion zone. This estimate is .lot very accurate because not all

of the oxide in the depletion zone Is dissolved. From inspection of

Figure 13 and similar photographs it was estimated that 75 percent of

the oxide in the depleted zone was dissolved. The calculation Indicated

o that the solubility of oxygen in SmCor at '125 C in excess of the solu-

bility at SOc/'c is _ 0,35 wt%. It is impossible to put error limits

on this estimate.

Next, the oxide content within the core of ar1 aged specimen was

calculated metallographically. Large oxide particles are produced as a

result of the aging treatment and these oxide particles appear black under

reflected light; see Figure 16. Consequently, a Quantimet image analysis

was performed on several sections and the ar;a percent of oxide was

determinrd, The average area of oxide particles and the limits were

determined to be 2.6 ± 0.4%, The area percent is equtl to the volume

percent for randomly distributed equidimensional particles in a dilute

suspension. Hence, from the volume of oxide in the core obtained after

aging, the excess solubility of oxygen at the sintering temperature was

determined. This calculation involves the implicit assumption that all

of the soluble oxygen in excess of the 800 C solubility limit was pre-

cipitated as measurably large oxide particles during the 45-minute aging o

treatment at 800 C. The precipitation and growth of extremely large

oxide particles during this period, as shown in Figure 16, indicates that

32

k'-A ^^^^ .

■M ----■■■-■-^- -:-.-. »w«..

-rr^MM^.^y.x-nmmm-lvmmmm ll i.il.il I ■ .1 i iw^^rwf«. i ■■lllwni^^n^n^P« mmmmmm w*m^mm wmmmmmmmnmQ

SA Rm 91

FIGURE 16 OXIDE PRECIPITATES IN THE SmCo5 GRAINS AFTER OXYGEN SATURATION AT 1125"C AND AGING AT 800 C

■Mi

kjt~am ■ l_llMaKaM —^^tmm jjt

--.■-IT'WVW; I,IIIUJIJ,I|IIIUI»IIMII' 'wi:f..iiiMijiiiu«!W1wimiii j I .lllinjFnmpmin>«W1(l,4)l>|i « IIPI|I(.1|J|IIJJ,P1W,I|WII| .UJlllW'fWmilll.HIHIWilllnlJIM wwm^l'aTO W-MFWI u I , •!, .WWBWBTOI

this assumption is valid. The metallographic determination of the o

excess solubility at 1125 C is 0.35 ± 0.05 Wt% (3500 + 500 ppm).

Although the oxide particles observed in the aged specimens were

sometimes larger than 100 micrometers, particles in the quenched specimens

are generally smaller than one micrometer. Any exceptions can usually

be traced back to larger oxide inclusions that were present alter melting.

The oxide particle population is greater in the quenched specimens than

in the aged specimens, as shown by cont.-asting the quenched and aged

sepcimens illustrated in Figure 17. Again, the specimens were castings

subsequently oxidized and annealed at 1125 before quenching. These

micrographs were obtained from surfaces polished using cerium oxide

but not etched. The background also shows a faint relief of larger

ym Co precipitates, which will be discussed in the next section.

However, a large amount of ^ygen known to be present in the specimen

is not accounted for xn the low particle population regions of the

quenched specimen, which is most of the specimen. This oxygen ray occur

as small particles below the resolution limit, as coherent oxygen-rich

zones, or as a supersaturated soluto. The aging process collects the

smallor oxide particles into a few large particles by a solution-

reprecipitation process.

The quenched specimen micrograph of Figure 17 shows a region of

low oxide particle population in the upper part of the figure and a

region of high particle population in the lower half of the figure.

High particle population regions were less common than low particle

population regions in the specimens examined. The boxed area in Figure

17 was picked at random within the high particle population region.

The average oxide particle size is 0.4 micrometer and tho volume frac-

tion of oxide, V , was determined from

Vf = J N d f 6 s p (3)

34

MM ^^Mfcl».! .. Mi . i

—■ ,.. ...,1.,.,. [ jiuiu^.iuji, uji>'n><*pf*iiiai<iOTnmmi|i.>iwiiiiii.ii«inw<ni«w''*>iuiii' i »i nimmawm^mmfimimmr^mmiwmmmil'^imimmiim'wnmKmifi "wm

v. -

» i

■

*

• *

.. ^v

■/ V

•

■.»..

* •

. F« * 20/J m

(iit QUENCHED FROM 1125 C

_ *

1

^ >r

, **"* ,

Ibl QUENCHED AND AGED AT 800 C SA HI \\ 90

FIGURE 17 QUENCHED AND AGED SmCo6 GRAINS

35

^^MMM.

wimi^Hmiiii. ■ i«HT^v^^iinp mwwwi^FtPH^Brwpwim^piw^ -»-^"^■MW

where N is the particle population per unit area, and d is the mean s P

particle diameter. The particle population in the boxed area was 7 2 3.0 x 10 particles/cm . The resulting volume fraction wus 0.025 (2.5%),

which is very similar to the volume fraction (2.6%) in aged specimens.

Hence, in the high population regions, most of the dissolved oxygen has

precipitated as submicron but visible particles.

The appearance of small oxide inclusions within the primary SmCor

grains of both experimental and commercial sintered magnets has also been

observed. These magnets did not receive an aging treatment. Small rxidc

particles generally appear within the SmCor grains rather than at grain

boundaries. Fracture surfaces have also been examined with a scanning

electron micrograph and oxide particles are rarely observed until the

specimen is etched. Numerous oxide particles which were present in the

SmCo grains are uncovered by etching. Exposed oxide particles in an 5

etched sintered magnet are shown in Figure 18. The dark areas in this

scanning electron micrograph are pores.

~m*m

"■■ — m1 »■

SA 871t 97

FIGUF^E 18 SAMAf^UM OXIDE PARTICLES IN A SINTERED SmCOg MAGNET UNCOVERED BY ETCHING

M

——-

V PRECIPITATION OF Sm Co IN SmCo 2 17 5

A study was made of the precipitation of the Sm Co phase In 2 17

samarium-depleted SmCor USIIIK cast pieces. There are two stages of

Sm2C"l7 l,ruciPitution ln SmCo as a result of partial oxidation followed

by sinterinj; or vacuum anneal in«. The primary staue results from samarium

depletion caused by reaction between the SmCo core and fi-cobalt In the 5

subscale. This reaction occurs durin« the sinter-anneal inn process de-

scribed by liquation (1). As a single phase alloy is quenched from 1125*^,

precipitates of Sm Co occur within the SmCo. «rain as described by

Equation (2). These precipitates are one micrometer wide and u few

micrometers Long. They often occur in clusters as shown by the ■canning

electron micrograph in Figure 19.

One small samarium oxide precipitate also appears in this mlcroKiaph,

The average SmCo5 «rain size, which is controlled by solidification, was

about 80 micrometers i- ')ur tcjt specimens. There wat no tendency 'or

Sm2Coi7 to PreclPltatc ut thc SmCo «rain boundaries during the primary

sta^e of Sm2Coi7 precipitation after vacuum annealing at sintering tempera-

tures. However, cobalt-rich specimens quenched from the melt exhibit

Sm Co at boundaries of SmCo grains. ^17 5

It should be emphasized that primary precipitation of Sm Co within

the SmCo5 grains occurs whenever excess cobalt is exsolved from a samarium-

depleted SmCo5 phase during cooling, whether the cause of samarium

depletion Is oxidation or not.

Primary Sn^Co^ precipitation is evidently complete on quenching

in vacuo. A series of aging tests for various times ut 800OC using pieces

from a samarium-depleted double molted ingot (not oxidized) showed no

increase in the amount of Sm Co or any apparent rearrangement of Sm Co tu 2 17

as a result of aging.

Preceding page blank

39

^*^k*m

SA 8731 98

FIGURE 19 SCANNING ELECTROMICROGRAPH OF Sm?Co,7 PRECIPITATE': IN SmCo&

Id

■HMMHH ^mm

1 ■■■II«

The second stage of Sm Co precipitation is caused by rejection

of oxygen dissolved in SmCo and formation of Ba 0 . As already mentioned,

this does not occur to any appreciuble extent during quenching, but It

does occur with aging. During the aging process, both samarium and oxygen

are collected at the large oxide particles and must diffuse over distances

relatively large when compared with the primary SmCor grain size. Further

depletion of samarium from primary grains and rejection of additional

Sm Co results. Micrographs indicate that this secondary precipitation 2 17

of Sm Co occurs at the SmCoE grain boundaries rather than within the 2 17 0

grain i. This is particularly evident in the aged specimcMi micrograph of

Figure 17.

The precipitation of large Sm 0 particles is probably controlled by

rapid grain boundary diffusion of samarium. This process causes bulk

depletion of samarium from the region adjacent to SniCor grain boundaries

and precipitation of Sm^fo in the samarium-depleted zone, which is about

2 microm^ers wide. riix« model is consistent with the r i c ol .-lid state

sintering of SmCo=, a process t mt is also controlled by grain boundary 5

diffusion.

Similar behavior has been observed in sintered magnets. However,

when the primary SmCo grain size is small, ~ 5 micrometers, the Sn^Co^

precipitates are no smaller than obsc.-'ed in cast pieces, and, hence,

nearly as large as the primary grains.

The precipitation of fine Sm Co particles in samarium-depleted

alloys contrasts sharply with the Sm Co phase in samarium-rich alloys.

In the latter case the primary SmCo grains are clean and Bm Co occurs

us separate grains outside the SmCo primary grains.

■11

^^^ • - -- ... ■ mm iaMu>^~aMM

v^rmmammm*

VI DISCUSSION OF Sm 0, A.M) Sm Q) INCLUSIONS ANL)

INTHINSIC COERC1VITY IN SINTEHKI) MAGNETS

A. Summary ol Inclusiona

It is now quite clour thut internal surfaces can occur in primary

Sm'"« Rrains of sintered magnets from precipitation of second phase 5

inclusions: Sm Co and SmrO,■ Although the oxygen solubility dotermina-

tions for cast SmCo arc not very precise, it is certain that large 5

amounts of oxygen can dissolve in SmCo. at sintering temperatures. For 5

careful processing operations, where oxygen contamination is kept below

3500 ppm, all of the oxygen will be dissolved in SmCo during sintering. 5

The normal quenching of sintered magnets is expected to leave approximately

one-micrometer oxide particles within the SmCor grains, which has been

verified in sintered magnets. Much of the oxygen in quenched SmCo is

uimccountod for and may bi prc^tnt as oupci turned oxygen, joheie.it

o;,ygen-rich zones, and unresolvablc oxide inclusions. Selective oxida-

tion of the alloy also affects the formation of Sm Co by depleting

samarium and fhifting the composition, and this occurs in two sequential

Bt age n.

Although oxidation cannot be entirely prevented, primary and secondary

formation of Sm Co precipitates can be prevented by having excess sama-

rium present in the powder compact to replace the samarium depleted from

SmCo grains. Oxide precipitates can be removed from the primary magnet 5

grains by an aging treatment that tends to collect the oxygen in a few

large oxide grains. Optimum aging conditions need to bo determined.

D. Inclusions and Intrinsic Coercivity

This section begins with a brief review of coercivity models for high

Preceding page blank 43

■

——-"•^PBP

magnetocrystalline unlsotropy magnet materials. For a detailed review,

see LlvitiKston. The highest published i-oom temperature intrinsic

12 coercivity in sintered SmCo. magnets is IL,. = 43 kOe. A sintered

5 1

SmCo magnet with a room temperature intrinsic coercivity of 43 kOe has 5 3

also been produced and measured at SRI. This is only about 15% of the

theoretical coercivity. 2 K/M , where K is the magnetocrystalllne anisot- s

ropy and M is the saturation magnetization. Furthermore, sintered SmCor s J

magnets usually have much lower coerclvitles and SmCor powders and

samarium-depleted sintered magnets usually have intrinsic coerclvitles

well below 20 kOe.

Because of the high mannutocrystalline anisotropy, coherent rotation

<JI domains and curling cannot occur at the comparatively low fields at

which magnetic reversal actually occurs.

Magnetic domain nucleation and domain-wall pinning arc factors that

may control coercivity. Inclusions and possibly other defects such as

stacking faults and coherent zones o! second phase material are expected

to impede domain-wall movement aid Increase OOorolvity« However, thes.e

same defects can act as nucleation sites for new domains and decrease

coercivity. The experimental results show conclusively that processing

operations that lead to Sm,0, and Sm Co inclusions within SmCo grain

boundaries always lower coercivity and convorsely processing operations

thai decrease these inclusions raise coercivity. Hence, nucleation of

domains rather than domain-wall pinning evidently controls coercivity.

Additional arguments against a domain-wall pinning model can be

made based on other e^perimontal information. Sintered rare earth cobalt

magnets are made by aligning the powder (easy axis crystallographlc

parallel orientation). During sintering the crystallographlc orienta-

tion is maintained but the specimen is thermally demagnetized. If

domain-wall motion controlled the coercivity, then magnetization of the

sintered specimen should require a high field similar to H . The ci

M

■^MM MM - . ,__

experimental evidence is that magnetization nearly equal to M can be

obtained in a thermally demagnetized specimen at fields much lower

than H . ci

Domains have been observed in a large number of SmCo^ grains in which 5

Sm Co inclusions have been introduced by samarium depletion. Bending 2 17 13

of domain walls around these inclusions or termination of domains at

internal inclusions has never been observed in this laboratory. In all

cases the domain walls seem to be oblivious to the presence of the cobalt-

rich inclusions, A typical micrograph of domains in SmCo grains with 5

Sm, Co precipitates is shown in Figure 20. A small domain apparently

growing out of the grain boundary which is a region of Sm Co can be

seen in this micrograph.

Small oxide particles, ^1 micrometer, such as those shovn in the

quenched specimen of Figure 17 also do not appear to have a preferred

locus coincidental with domain walls and, therefore, are evidently not

pinning aom '..i wuiii. However, we do at' ;ia"L much evlllfenoc on t i , yet.

An inclusion domain-nucleation model can explain most if not all

of the confirmed phenomenological relations between samarium-cobalt

magnet processing and the resulting intrinsic coercivity. By confirmed

phcnomcnological relationships we mean behavior that has been observed

in more than one laboratory or is generally accepted. An example is the

observed drastic decrease in coercivity if excess samarium is not present

during sintering. According to this model, both Sm 0 and Sm Co

inclusions, which may be submicroscopic, can heterogeneously nucleate

new domains. Magnetic coercivity of a grain and of a sintered aggregate

will decrease as the population of small inclusions increases.

Rather than present a lengthy discussion of these phenomena and

their possible association with inclusions, we have chosen to list effects

with postulated causes in Table 1.

45

•-^^mrmm^m^rm^^m <*m^mmmi^^^™9^^mfm^^mmmri^^^mt " "\*\"^^m^m^^m^*^m*:^qfi^w*i

SA R731 99

FIGURE 20 MAGNETIC DOMAINS PASSING THROUGH Sm^Co^ PRECIPITATES WITHIN SmCo5 GRAINS

46

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47

mmm—•—— ■ ■Illl 'Ill "---

< ii.iiwiwwiui.LJii^mvji^wfijp^ii!JIIU|W^^^

VII SINTERING

A. Introduction

Benz and Martin studied the mechanism of sintering of rare-earth

14 cobalt alloys, and postulated that the sintering kinetics are controlled

by the diffusion of samarium atoms in grain boundaries via a samarium

atom-cobalt vacancy cluster exchange mechanism. This conclusion was

based on an observed one-third time dependence of volume shrinkage

calculated f .'om end point densities of samples that initially contained

a liquid phase, and on a comparison of the total volume shrinkage or

density with the alloy composition. Benz and Martin further noted that

a correlation existed between the conditions that produced rapid shrinkage

and the conditions that yielded high intrinsic coercivity, and since

this correlation has important processing implications, it was decided

to follow the solid phase st^ge of sintering of SmCo_ under conditions o

in which kinetic models could be applied thus alloving a study of the

effect of changes in stoichiometry on the kinetics of sintering.

The sintering kinetics of solids have been studied extensively 15 16 17

beginning with Kuczynski and then followed by Kingery and Berg, Coble,

and Johnson and Cutler. 18

Refinements in the assumptions relating the

mathematical models to the physical model have been made primarily by

19 Johnson, From these studies the general equation for solid state

sintering is given as

Preceding page blank

AL

L

Kvno

kTr

m

(4)

49

^.^tttmimm

■UPlilli I iuiwuii ^iJHJi||MHfH!JHI|tiUiii!iy||Llpujj|JI|!JliPili|Jii ^pffl mmmmwmimmiwvm. m u • ■ ü 11 i i

where

AL/L = the fractional shrinkage o

V = the surface free energy

Q = the volume transported per atom of the rate-controlling species

D = the self-diffusion coefficient of the rate-controlling species

k = Boltzmann's constant

T = absolute temperature

r = the particle or powder radius

K, n, m = constants that depend on the sintering mechanism.

Table 2 lists the values of the constants versus the sintering

mechanism.

Table 2

SINTERING CONSTANTS FOR EQUATION (4)

Sintering Mechanism K n m Reference

Volume Diffusion 5.34 3 0.49 5

20/V2 3 0.40 2

2 3 0.50 3

Grain Boundary Diffusion 2.14b 4 0,33 5

3b 4 0.33 3

The symbol b in Table 2 is the grain boundary width. The constant m is

usually determined by plotting the logarithm of fractional shrinkage

versus the logarithm of time, and a sintering mechanism is deduced based

on the value of m. This method of determining the sintering mechanism

50

mmmmtm ite — i^M ~—'- ■

Is not precise unloss time und IcnKth corrections are applied, because

m varies with small errors in time and shrinkuKe. The errors result

from the inability to heat the sample instantaneously to the sintcrinn

20 temperature. Lr.y and Carter have outlined a relatively ea.sy method

of correctin« the time and 'onuth that involves calculating L from

thermal expansion data and extrapolating a length versus time plot to

Intersect the calculated L value, thus determining the effective initial 0

t ime,

21 Johnson later pointed out that the method of Lay and Carter is

strictly valid only if an array of identically sized spheres are used in

the study, and if it is known that oil of the mass transport is due to

u single mechanism.

The sintering studies described in this report are incomplete in

that a complete range of itoichiometries has not been investigated and

the validity of assuming u single transport mechanism has not been tested.

D. Experimental Procedure

Samurium-cobalt alloy powders were made for this study by arc

melting the elements, Slight changes in the alloy compositions were

obtained by adjusting the samarium/cobalt ratio used in the arc melting.

The arc-melted alloys were annealed in an oxygon-gettered argon atmosphere o

at 1110 C for four hours, and were then individually ground in an alumina

ball mill for three hours using sodium-gettcred hcxane as the milling

fluid. The final alloy compositions were determined from EOTA titration 22

data, ' and the compositions were adjusted for loss of samarium due to

the formation of Sm 0, . The oxygen concentrations in the alloys were

determined by neutron activation analysis and by vacuum fusion analysis

(courtesy of General Electric Company). The oxygon content in solid

solution was assumed to be zero.

51

^■^M

Disk-shaped specimens 1.0 inch in diameter by U.13 inch thick,

were fabricated by die pressing at 30,000 psi in the absence of a mag-

netic field. The specimens were placed on a molybdenum support in a

furnace employing a graphite resistance heater surrounded by oxygen-

gettered argon. The change in sample diameter was recorded by time-lapse

photography as functions of time and temperuture, and the data wore

read fron the films by means of a Teleroadex film analyzer.

Sample length corrections were made following the method of Lay

20 and Carter using the thermal expansion data presented in Figure 21,

which were obtained by using a commercial dilatometer modified so that

a static argon atmosphere could be used to surround the sample to prevent

oxidation.

Heating the sintering snmples from room tempi'raturc to the sintering

temperature requiroü approximately 90 seconds. Shrinkage below approxi-

o o mately 1100 C is almost negligible and the time required to reach 1165 C

from 1100 C was approximately 5 seconds. Calculated L values (initial o

sample diameter) corresponded exactly with the experimentally measured

o o L values at 1100 C. At 1120 C a small correction in L was required

0 o and these corrections increased in magnitude as the sintering temperature

o was raised to 1165 C, indicating that shrinkage was occurring as the

o sample was heated from 1100 to 1165 C. The corrections ranged from 0.0%

at 1100OC to less than 2.0% at 11650C.

C. Results und Discussion

Shrinkage data were obtained for four different alloy comparisons,

and these data are shown In Figures 22, 23, and 24. The value of m

(sec Eq. 4) as determined by a least-squares analysis ranged between

0.24 and 0.37. The correlation coefficient except for one isotherm was

above 0.93. Thus the shrinkage data approximate a slope of m = 0.33 and

indicate that the shrinkage is controlled by grain boundary diffusion.

52

i^MMiM —

iMWHM

1 4 1 1 1 1 1 1 1 1 >

1 2

§

'—

, ,0 — y

EX

PA

NS

ION

o

as — v^

<

-

< 06

■ Ul z

yr *- 04

< a Z y* ' 3 0.2

0 sS 1 1 1 1 1 J__ 1 1 1 100 200 300 400 5UU 600

TEMPERATUBE - 700 800 900 1000 1100

SA 8/31 103

FIGURE 21 LINEAR THERMAL EXPANSION OF lJOLYCRYSTALLINE RANDOMLY ORIENTED SmCoc

53

■MB^Maa IIHIHBIII ■ h^< ^■M - -*- ^■■■~ -'^ JM

^^

8 8

I I z

\

cc O u. in 5 CC

8 HI o < z cc z (/)

< 00 o o T s £ < s < CO

CM CM

CC

o d

0-inv

8 d

54

■ ■ ■ 1111 m liiiiiiinin r

^^^w^^^

o

<

I O

UJ

I i i

S

-inv

55

MaM MütfiüitaiaMaii^^kM^MM^MMiMiM •aMMMMMM

■ IL-I 1 ' •^m^^

8 S

E 1/5

n CO

n Q 2 <

CO <■>

a z z <

o u to > o

cr O LL

(/) s tr LU

g

X 10

3

UJ

nnv

56

MMM ii »iiii—MiMMinriii ,

^^»-v^Mm^ ■■ - •^m

The data presented in FiRUres 22, 23, and 24 were force fitted to lines o

drawn with slopes equal to 0.33, and the data near 1100 C exhibit con-

siderable scatter due to the small amounts of shrinkaRO involved. The

rate of shrinkage can be calculated from the intercept values for each

isotherm, and the rate of shrinkage is shown plotted versus reciprocal

temperaturos in Figure 25, All four alloy compositions exhibit identical

rates ol shrinkage within experimental error and this result is in conflict 14

with the results of Benz and Martin.

SmCo alloy compositions are very sensitive to the oxygen impurity 5

concentration. Therefore thr oxygen content of the alloys used in this

study were determined by both neutron activation analysis and by vacuum

lusion analysis. These methods of analysis for oxygen in SmCo give

conflicting results with the vacuum fusion method always exhibiting the

higher value. The results of the oxygen analysis are listed in Table 3.

Because of the discrepancies shown in Table 3, very brief descriptions

of the two oxygen analysis techniques are given below.

The neutron activation analysis was conducted by Gulf Radiation

Technology Company. Each sample was irradiated for 10 seconds in a 14-MeV 8 2

neutron flux of approximately 10 neutrons/cm -sec. An oxygen standard 16

was irradiated with each sample and the induced activity of N was

counted on a single channel pulse-height analyzer using a pair of

5 x 5-inch Nal(Tl) scintillation crystals. The N is formed by oxygen 16 16

interacting with neutrons according to the reaction 0 (n,p)N . The

oxygen concentration was determined by comparing the intensity of the 16

6.13-MeV gamma-ray photo peak of N from the sample with the same peak

from the oxygen standard.

The vacuum fusion analysis was conducted by General Electric Company

and an excellent description of this technique has been given by Horten

57

MM t^mm mmmmmm*immm .M^MMH .

•^wmmmmmm mmmmm **rm

0.10

0.09

0.08

0.07

0.06

0.05 -

0.04 -

0.03 -

i r

^"y 0.02 -

0.010

0.007

• -27.7 w/o Sm

■ 31.4 w/o Sm

A 33.0 w/o Sm

A 33.3 w/o Sm

1/T x 10

FIGURE 25 ARRHENIUS PLOT OF THE SHRINKAGE RATE VERSUS RECIPROCAL TEMPERATURE

58

mmmm ^^^aMMMMi

■ II II I III! PI

■

/^ t' b » ■ >^ 0

>> i m 0 5 rH n pH

pH • • t

(—4 pH 0) s 1 < E M M

3 c 3 •^ U

a c > 0 •H

a u *i

G c ■ 0 0 •H e 4->

8 a >

•H m s tf CO c H t*

w Ü • • • • < n n rH tv

•o n n co N • c •p p M K

■H 4-> 3 3

CO U V >> pH 2 O ■ a a o

w 1 ~ <

m c 0 0 o •H E tn

M m 5 *

m u. U • • •

01 O 1 g pH pH o 1

PH ^/ 3 Si w 3 a M c B H

1 0

■H

cs

(4 >

3 < e

0 •rl

% B il H D a

e > g 8 m «0 «f CO Q U (0 O) * CO

c < 00 in m 0)

V • « • • bC e o o o o >> 0 R (H

o -P 3 i c 0 •rl

>. P A U

C9 t* t*

m -o ^ PH p-l

a> u ■H in m m o f 8 «1 -P "H CS U -H 8 8 U M .; o Q B e E E E E a +J w m CO CO W CO

0) >> a d «V —*>»

H 1

X

>> f- 00 OJ o 0 • CO «5 <c t*

pH 0 1 1 1 1 PH K s s ■ S <

59

— —^

" " «■...- —^-. .-..-...- . -.-™L. - . J I.,. ,..„,.. L I RMIIM

23 and Cnrson. The technique cssentJally consists of melting the SmCto

D

alloy in a platinum-tin bath in a vacuum in the presence of an excess of

carbon. Carbon monoxide is produced and the amount produced is taken as

a measure of the oxygen content in the alloy. The alloy sample is placed

in a platinum foil-tin container and introduced into a vacuum side arm

above an outgassed graphite crucible containing 95% platinum plus 5% tin

o -5 held at 900 C in a vacuum of approximately 10 t^rr. The sample is

dropped onto the solid ilatinum bath and the tempen'ture is raised to

o o 1950 C requiring approximately 12 minutes to reach 17.>0 C and approxi-

o mately 15 to 18 minutes to reach 1950 C. The quantity of gas evolved is

determined manometrically using McLeod gages.

Oxidation of powder samples or porous samples may occur during the

time the sample resides on top of the platinum bath and this might account

at least in part for the higher oxygen values obtained when analyzing

by a vacuum fusion technique.

Table 3 also lists the phases detectable by x-ray diffraction. If

the oxygen contents determined by vacuum fusion analysis are correct, we

would expect that alloy M-68 would contain a greater amount of Sm Co

than the M-69 alloy. Since this was not confirmed by x-ray analysis,

all of the alloys were examined metallographically. M-67 and M-68 con-

tained very small amounts of second phase and M-69 contained much more

second phase than did M-68, but less than M-70. Therefore, the x-ray data

and metallographic examinations correlate more closely with the neutron

activation data. It should be realized that the sintered specimens sub-

mitted for oxygen analysis were very porous because only the initial

stage of the solid phase sintering process was studied, and the possibility

of oxidation occurring during the vacuum fusion analysis does exist.

Additional work is necessary to determine the best method of oxygen

analysis in SmCo alloys; however, in this report the neutron activation 5

analysis data have been used to calculate the samarium contents of the

60

.—^^^^^^^ . .__. i—m ii—im i : üäi

■ •-■""I wmmnmimmmmwwi'iim>inmpmm «i.ii"infim**m~***mmmmmm**™**^v*»*u***v™ ^mm^m^mmm

alloys (see Figure 25). The samarium content has therefore been varied

almost entirely across the SmCo phase field and into the Sm2Co

17 Phase

field with the exception of the postulated hyperstoichiometric SmCo

24 region. Since we do not observe any difference in the ratio of

shrinkage as the samarium/cobalt ratio is varied, we are forced to con-

clude that deviations in stoichiometry in this system do not alter the

grain boundary defect structure sufficiently to cause a change in the

grain boundary transport of the rate-determining species.

Additional work will be completed during the next report period

that will allow a quantitative evaluation of the solid phase shrinkage

rate, including experiments to extend the samarium/cobalt ratio into

the Sm Co phase field. This work is necessary before the final con- 2 7

elusions from the solid phase sintering work are made.

61

mmmmt^mm «a^mfm tmmmm

|!IWI<|I|II'I,'III<I'I"<11111 immmmmmmmmmmmm*** r**mm mmum ^mmmmmmmmn

REFERENCES

P. Kofstad, Nonstolchiometry, Dlifusion, and Electrical Conduc-

tivity in Binary Metal Oxides, Wiley-Interscience, p. 28G,

New York, 1972

2. T. L. Felmlee and L. Eyring, J. Inorg. Chem., 1, 660-66 (1968)

3. P. J. Jorgensen and R, W. Bartlett, "Materials Processing of

Rare Earth Cobalt Permanent Magnets," Tech, Report AFML-TR-72-225

(September 1972)

5.

C. E. Wicks, and F, E. Block, Thermodynamic Properties of 65

Elements - Their Oxides, Halides, Carbides, and Nitrides,"

U.S. Bureau of Mines, Bull. 605, Govt. Printing Office (1963)

P. J. Jorgensen and R. W. Bartlett, "Materials Processing of

Rare Earth Cobalt Permanent Magnets," Tech. Report AFML-TR-72-47

(April 1972)

6. G. D. Stone, Ph.D Thesis, Arizona State University, 1968

7. C. D. Wirkus, M. F. Berard, and D. R. Wilder, J. Am. Ceram. Soc.,

50, 113 (1967)

8. M. F. Berard, C. D. Wirkus, and D. R. Wilder, J. Am. Ceram. Soc,

51, 643 (1968)

9. V, B. Tare, and H, Schmalzried, Z. Phys. Chem. N.F., 43, 30 (1964)

10. M. G. Benz and D. L. Martin, J. Appl. Phys., 42, 2786, (1971)

11. J. D. Livingston, "Present Understanding of Coercivity in Cobalt-

Rare-Earths, to be published in AIP Conference Proceedings (1972)

12. K. H. J. Buschow, P. A. Naastcpad, and F. F. Westendorp, IEEE

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13. H. Zijlstra, J. Appl. Phys. 41, 4881 (1970)

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