Influence of HDDR treatment on magnetic properties of TbFe2- and DyFe2-based intermetallics

AND COMPOUNDS ELSEVIER

Journal of Alloys and Compounds 260 (1997) 271-276

Influence of HDDR treatment on magnetic properties of TbFe 2- and DyFez-based intermetallics N.K. Zajkov*, N.V. Mushnikov Institute of Metal Physics, 620219 Ekatr ,nburg, Russia

Received 7 March 1997; accepted 27 March 1997

Abstract Magnetic properties of TbFe 2, DyFe 2, Tb(Feo.sMo.2) 2 and Dy(Feo.aMo.2) 2 with M=Co, AI, Si, Ga alloys affected by the Hydrogenation-Decomposifion-Desorption-Recombination processing have been studied. After hydrogen treatment the coercive force H~ grows sharply, so HDDR-powders can be used as isotropic permanent magnets with the energy product up to 26 MG Oe at T=77 K. In Dy(Feo.sAlo.2)2 at T
1. I n t r ~ u c t i o n Hydrogenation-Decomposition-Desorption-RecombP nation (HDDR) processing is widely used to realize a small-grain-sized structure in Nd2Fe,4g, SmC%, Sm2Co~TN ~ and some other alloys for permanent magnets ttj."" .,u~,~'--"~ t.,~. ,.". the. . . . HDDR treatment at optimum conditions, such compounds show a coercive force (He) comparable with that of sintered magnets prepared by the usual metalloceramic technology [2]. However, as a rule, HDDR-powders have a random orientation of the easy axes of crystallRes and, consequently, are used for me production of isotropic magnets (in particular, bonded magnets). In this case the remanent induction B~ of such an ensemble of randomly oriented umaxial particles does not exceed 0.5 B~ (where B~ is the saturation induction). Preparation of anisotropic HDDR-l~wders is a traditional way to obtain B~ increase. This can be achieved either by the alloying of such metals as Co, Zr, Ga to the basal composition [3] or by applying a non-equilibrium I-IDDR process [4]. At the same time, the ensemble of randomly oriented paaicles with a cubic at~sotropy is known to have essentially a large B~/P~ value. In pa~icular, for a < 111 > type easy axis polycrystalline sample, BJB~=0.866. Thus, *Corresponding author. 0925-83881971517.00 © 1997 Elsevier Science S.A. AU rights reserved. Pll S0925-8388(97)00168-0

cubic materials are perspective for isotropic permanent magnets from the point of view of enhanced remanence. The TbFe 2 and DyFe 2 compounds possess a MgCu 2type f.c.c, crystal lattice (Fd3m), high values of the Curie temperature Tc, saturation magnetization and magnetocrystalline anisotropy constants. However, a low coercive force prevents their application as permanent magnet materials. A partial substitution of Fe by Ga or AI increases the coeicive force at low temperatures [5,6], which allows the use of these compositions as cryogenic permanent magnets. As was shown earlier, in quasi-binary alloys Dy(Fe,M) 2 (M=Ga, AI, Si) the stepwise magnetic reversal is observed at low temperatures [7-12]. This effect was attributed to different reasons: a stepwise redis~bution of magnetic phases in the sample [7], pinning of narrow domain walls [10,11], effects of quantum tunneling [12,13], effects of domain wall "inertia" in a cer~in critical ~[a..~ _ ~._j.r191We have shown [8] that the magnetization jumps in the given alloys are caused by a sharp change of temperature in a sample as a result of heat released by the irreversible motion of narrow domain walls. In this work the influence of HDDR treatment on magnetic hysteresis properties and the effect of low-temperature stepwise magnetic reversal of TbFe 2- and DyFe2based interrnetallics has been repotted.

272

N.K. Zajkov. N.V. Mushnikov I Journal of Alloys and Compounds 260 (1997) 271-276

The compounds investigated were prepared by the melting of initial components in the induction furnace in an argon atmosphere with the subsequent annealing at 1000 °C for 24 h. According to X-ray diffraction data, the binary compounds TbFe 2 and DyFe 2 contained less than 3% of extraneous phases. The thermomagnetic analysis shows that quasi-binary compounds Tb(Fe,M) 2 and Dy(Fe,M)2, with M=Co, AI, Ga were practically singlephase tco. In Tb(Fe,Si) 2 and Dy(Fe,Si) 2 impurities of extraneous phases of about 10% were observed. The HDDR treatment of the investigated alloys was performed using the following standard procedure. Directly before hydrogenation the sample surface was activated by heating up to 200 °C in vacuum 10 -5 Torr. Then the sample was exposed to pure gaseous hydrogen obtained by the decomposition of LaNisH x hydride. After a complete hydrogenation, detected by the termination of pressure change in the working chamber, the sample was heated up to 660-820 °C and kept at fixed temperature and hydrogen pressure 1 bar for l h. This provides the decomposition of initially formed RFezH,, hydride into RH 2 hydride ~ d a-Fe. Then hydrogen was extracted from the sample in dynamic vacuum, and the initial compound was recombined. It should be noted that the time needed for a complete desorption of hydrogen from samples determined by the restoration of the magnetization, for DyFe, is approximately three times longer than for TbFe,. After complete aesorption of hydrogen, the sampler, were cooled down to room temperature. The samples after HDDR treatment become brittle and spontaneously disintegrate into a powder with the mean particle size 30-70 p,m. For magnetic measurements HDDR-powders were pressed under 5 10" Pa pressure into pellcts of 4 mm in diameter and !.5-2 mm thick. The magnetic hysteresis properties of samples were studied on a vibrating sample magnetometer in magnetic fields up to 20 [:De. The Curie temperatures of ferromagnetic phases were determined from measurements of the initial magnetic susceptibility.

3. Results and diseussio~i

3. i. Hysteresis magnetic properties

The room temperature coercive force dependence of TbFe 2 HDDR-samples on the temperature of hydrogen treatment is shown in Fig. 1. As can be seen, after HDDR treatment H c grows by an order of m gnitude in comparison with the initial compound (H~ = 130 Oe) and has a maxinium at the treatment temperature T,r=700°C. At lower temperatures a smaller He value is likely to be related to the incomplete recombination and/or the incomplete desorption of hydrogen from the sample. The latter can be indicated also by the reduced B r (Fig. I) and Tc

10

3.5

2. E x p e r i m e n t a l details

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2.5

.,===...-

m 2.0

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0

(.9 "4

t,_

1.0 0.5 0.0

I

6 0

!

i

700

750

.

tr

•

I

800

(oC)

Fig. 1. Room temperature dependence of the coercive force H~ (open circles) and remanent induction (closed circles) of a TbFe_, sampleon the temperature of HDDR treatment. values. At elevated Ttr '.he ='eduction of a coercive force can be attributed to the.: increase of the crystallite sizes owing to a recrystallization. Thus, there is a temperature range of HDDR treatment of the TbFe 2 compound (700-750 °C), where the H c value read : 2.5-2.7 kOe at room temperature. For DyFe 2 H c inc, eases from 100 Oe in the initial alloy up to 1.25 kOe in the HDDR-sample after treatment at 700 °C, and this va!ue weakly depends on T,r between 700 and 780 °C. The shape of an initial magnetization curve in both compounds is consistent with the domain wall pinning mechanism of He. This mechanism was previously used for the explanation of high H~ values at low temperatures in Dy(Fe,M) 2, where M=Ai, Ga ": '~ H,-,,u,=ver. the substantial contribution to H~ growth of HDDR treatment can be attributed to the nucleation process. As a matter of fact, as DyFe, and TbFe: possess a high magnetocrystalline anisotropy constants (the cubic Kj(0) constants are - 5 . 2 and 4.7 (× 108 erg cm-~), respectively [14]), the critical size of singledomain particles can be evaluated as 1-2 p,m. At the same time, the typical grain size in HDDR-materials does not exceed 1 p~m [15]. Fig. 2 shows the temperature dependencies of He, B r and the energy product (BH) m of HDDR-samples of TbFe: and DyFe: treated at optimum conditions. These materials can be used as cryogenic isotropic permanent magnets. As seen from Fig. 2a and 2b, at low temperatures DyFe 2 has the best magnetic properties, whereas at high temperatures TbFe 2 has. The values (BH) m are close, for example, to that for Dy(Fe,Ga)2-based magnets prepared by traditional metalloceramic techr~ology [5]. However, due to a less sharp decrease in H c with the temperature growth in HDDR-samples, they do not demagnetize by an internal field even at room temperature, whereas for Dy(Fe,Ga) 2 the coercive force has already decreased down to 1.5 kOe at 130 K [5]. In order to improve the magnetic hysteresis properties of HDDR-treated samples, we have prepared quasi-binary Tb(Feo.8Mo.~) 2 and Dy(Feo.aMo.2) 2 compounds with M = [J,uj.

=,... . . . . .

.

N.K. Zajkov, N.V. Mushnikov I Journal of Alloys and Compounds 260 (1997) 271-276

'10

~

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6 .,,,-,,.,

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100

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200

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300

n

!

400

T (K)

T (K)

c)

~" 20k 0 C9

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r (K)

3(10

400

Fig. 2. Temperature dependencies of Ihe coercive force (a), remanent induction (b), and maximum energy product (c) of TbFe 2 (closed circlesj and Dy ~t,.. ~:~2 (oWn circles) samples affected by HDDR treatment at 750 and 700 °C, respectively.

Co, AI, Si, Ga. Such substitution was c~qied out by the following reasons. The substitution of some of the Fe atoms by cobalt should result in Tc growth due to the increase of exchange interaction energy in a 3d-~ublattice. As seen from Table I, this is really observed for both types of alloys, however the room temperature magnetization value decreases. Substitution of Fe in TbFe z and DyFe 2 by

Table 1 Curie temperatures, magnetizations in a field of 20 kOe and coercive forces of Tb(Fe,M) 2 and Dy(Fe,M) 2 Compounds

TbFe 2 TbCFeo,Coo 2): Tb(Feo,sSio2)z Tb(Feo 8Alo ,)2 Tb(Feo 8Gao z) z DyFe, Dy(Feo sCoo 2) z Dy(FeoaSio 2): Dy(Feo ~AIo 2)2 Dy(Feo.sGao 2)~

7", (K)

720 760 570 445 450 628 673 455 365 523

i (kG)

H, (kOe)

-,'7 I," I i it,L

293 K

77 K

293 K

9.7 1 !.7 l 1.0 9.8 7.8 9.8 9.7 9.8 12.3 8.5

7.8 5.8 5.1 7.9 7.D 8.1 7.2 5.0 6.5 4.3

0.38 0.5 2.76 |.88 2.57 O. 1 0.16 0.15 0.09 1.00

O. 13 0.18 0.35 0.1 0.2 O. 1 0.1 0.15 0.05 0.1

other elements (AI, Si, Ga) should result in the following effects. First, as these metals do not have a magnetic moment, their introduction into a Fe sublattice shouid cause the ~'eduction of their magnetization and an increase of the total magnetization owing to fernmagnetic ordering of the magnetic moments of rare earth and iron sublattices. However, this effect will be accompanied by a 7",._. decrease, which can significantly reduce the effect of the magnetization growth, especially at high temperatures. As can be seen from Table 1, the increase of the magnetization is observed only for the compositions with AI and Si and only at low temperatures. The second expected effect from substitution should consist of a He increase [5,6]. In fact, the difference between ~he atomic radii of Fe and the substituted atoms will result in local distortions of a crystal lattice, that are displayed experimentally in a broadening of X-ray diffraction lines and result in a fluctuation of exchange interactions and local crystal fields [10]. As shown in Table 1, in practically all cases the substitution causes H c growth in compafisor, with parent TbFe 2 and DyFe z. HDDR processing of quasi-binary alloys also resuhs in a Hc increase. In Table 2 the optimum magnetic hysteresis properties of all investigated alloys after HDDR treatment

N.K. Zajkov, N.V. Mushnikov ! Journal of Alloys and Compounds 260 (1997) 271-276

274

Table 2 Optimum magnetic properties Dy(Fe,M) 2 at 77 and 293 K Compounds

TbFe 2 Tb(Feo.sCOo.2 ) 2

"I'o(Feo 8Sio.2)2 ~o(Feo.sAlo 2)2 Tb(Feo.sGao.2)2 DyFe 2 Dy(Feo.sCOo.2)2 Dy(Feo.sSio.2)2 Dy(Feo 8A1o.2)2 Dy(FeosGao.2)2

of

HDDR-treated Th(Fe,M) 2 and

10

(BH)m (MG Oe)

5

BHc (kOe)

B, (kG)

77K

293K

77K

293K

77 K

293K

4.0 1.2 2.5 3.6 3.8 8.2 4.0 4.7 3.9 4.2

2.4 0.5 1.3 0.9 1.1 1.2 1.0 0.7 0.5 0.6

8.3 5.5 3.8 6.0 6.5 10.6 8.6 7.0 1 i .9 5.8

6.2 4.1 2.5 3.1 4.3 7.0 5.0 2.4 3.0 2.2

13.0 3.0 3.5 7.0 8.1 26.4 11.3 9.5 22.0 7.2

5.4 0.2 1.2 1.1 1.9 4.3 1.6 0.5 0.7 0.6

(.9

0

m

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,

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are given. Table 2 shows that the substitution of Fe by other metals fails to increase the maximal energy product of HDDR-samples. A smaller amount of substituted metals (less than 13 at.%) can be expected to give a positive result. However, it is obvious that the essential improvement of magnetic properties in this case also cannot be achieved.

L9

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v

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3.2. Low-temperature magnetization jumps

;o

io

H (kOe)

In parent alloys Dy(Feo.sAlo.2)2 and Dy(Feo.8Sio.2)2 the stepwise magnetic reversal is observed at low temperatures (4.2-10 K), as shown in Fig. 3a for the compound with AI. This effect was studied in detail i:~ L8,9]. The magnetization jumps were shown to occur due to a heat release in samples with an avalanche movement of domain walls having a narrow distribution of critical fie;Js of pinning. At low temperatures when the specific heat is rathe,: small, this hysteresis heat can considerably increase the sample temperature. Fig. 4 shows "--m=temperature u~,u=,,~u~,-' . . . . -~. . . . . . u, H= of both initial and HDDR Dy(Feo.sAlo.2) 2 samples. In the temperature range when the magnetization jumps are observed (T< 10K), the definition of H= is rather complicated. For the HDDR sample it was carried out in an assumption that the H¢ value is approximately equal to the internal critical field of the jump (closed squares in Fig. 4), as shown by a dashed line in Fig. 3b. It is easy to estimate a heat release in a sample. When the magnetization increases on the valt:e A / i n an external critical field H , , , the amount of heat released during the jump is equal to work of magnetization and can be written

Fig. 3. The hysteresis loops for parent (polycrystalline sphere) (a) and HDDR-tre,ated at 700°C (pressed powder pellet) (b) Dy(Feo.sAlo.2)2 ~amplpc, ..me_~uredat -,.,~ ~*" K at the external magnetic field sweep rate 500 Oe s-~. The inclination of the dashed lines dlldH corresponds to the demagnetization factor of the samples.

can be determined from the equation: T2

Q= C (T)dT

(2)

T=

0

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e

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,

=

--"-'-

8~ ,...

"-"

0

g

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H.,,.

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This will increase in the sample temperature from T~ up to T 2, at which the H c value is lowered (see Fig. 4). Hence, thi~ state corresponds to the higher magnetization value in an external field H,i t, which is a driving force of the magnetization jump. The temperature of a sample heating

5 .

4

.

.

8

.

. . 12 T (K)

16

0 ~

Fig. ~ Temperatua~e dependencies of the coercive force for parent (open circles) and HDDR processed (closed circles) Dy(Feo.sAIo.2)2 samples, and the specific heat for DyAl 2 (solid line) I16]. Low-temperature H c values (closed rectangles) were detem'fined by the extrapolation of the critical fie~d of the jump on the ! = 0 axis (dashed line in Fig. 3b).

ILK. Zajkov, N.V. Mushnikov I Journal of Alloys and Compounds 260 (1997) 271-276

For a parent sample the magneti.zation jump AI = 826 G in a magnetic field H,~, = 11 kOe causes a heat release Q = 2 . 8 106 erg reel -~. According to Eq. (2) and Cp(T) dependence shown in Fig. 4 [16], this jump will heat the sample from 4.2 up to 15 K. A HDDR-sample Dy(Fe,AI)2 at A / = 8 2 6 G, H,~t= 18.4 kOe and Q=4.6 106 erg mol -~ will be heated from 4.2 up to 17.2 K. The thermal gauge attached directly to the sample allowed us to measure the temperature change during the stepwise magnetic reversal. The parent ~ample and the HDDR sample were found to heat up frt,:a 4.2 to 10.6 K and 11.1 K, respectively. After a sharp temperature growth, the samples cooled down to liquid helium temperature for 1-2 see. The difference between the calculated and the measured T2 values together with a fast cooling rate after the jump can be attributed to the non-adiabatic conditions of the experiment, since we tried to repeat the same thermal conditions as those at magnetization measurements. Besides, the expected value of the temperature jump on the HDDR-sample cannot be registered due to the relative low thermal conducfivRy caused by a weak thermal contact between particles of a powdered sample. As it was shown in [8], both the value and the number of magnetization jumps in initial samples depend on the sweep rate of the external magnetic field. At the sweep rate less than !0 Oe ~- ~ it is possible to obtain an isothermal hysteresis loop, without appreciable magnetization jumps (dashed line in Fig. 3a). For the same compositions after HDDR-treatment the effect of stepwise magnetic reversal becomes more distinct (Fig. 3b) and, besides, it is impossible to obtain a "smooth", isothermal demagnetization curve for any small sweep rate of ml external field below 7 K. To explain this difference let us consider the magnetic reversal process for a multidomain sample in a greater detail. When the magnetic field increases up to the lowest critical field of domain wall pirating Hc~,, a domain wall sta,~s moving to reduce the magnetic energy of a domain having the magnetization orientation opposite to floe applied field direction. This small magnetization change (the Barkhausen jump) is accompanied by the reduction of both the internal magnetic field and the field H,~, due to the local heating of the neighboring regions of the domain wall. If H , , decreases fast enough with the temperature growth, the domain will collapse. The thermal energy proportional to the domain volume, magnetization and H , , values will dissipate in the sample and heat other domain walls. But due to non-adiabatic conditions, a part of the thermal energy will come out from the sample. If the external field changes slowly, the sample can cool down to the starting temperature before the internal field reaches its critical value again, in this case the magnetic reversal will consist of small jumps in separate domains. If the external field increases rather fast, the sample heating can substantially decrease H¢~t, and this can initiate an avalanche domain wall motion, i.e., a complete or partial stepwis:

!

i

i '=

275

=

=

f

(.9 v v m

-8 -12 -20

-~8

- 6

-14

- 2

H (kOe)

Fig. 5. The staircase magnetization reversal for HDDR Dy(Feo,sAio.z) 2 sample at 6.9 K. The magnetic field sweep ra!e _equals 10 Oe s -~.

magnetic reversal. The same scheme is valid in Lhe case of single-domain particles in which the magnetic reversal is caused by the nucleation mechanism. The above described two variants of magnetic reversal are realized in HDDR-samples in different temperature ranges. Below 6 K, the stepwise magnetic reversal is observed at any slow field changes (Fig. 3b). At T > 10 K the demagnetization curve has usual smooth shape, whereas between 6 and !0 K at a small field sweep rate (10 Oe s -~) the demagnetization curve consists of clearly indicated small steps (Fig. 5). As the values of the magnetization jumps in this case are much larger than the magnetization of a separate domain, these jumps can be attributed to the remagnetization of either some domain groups or groups of powder particles having nearly equal Hcrit values.

4. Conclusion The cubic TbFe 2, DyFe 2 intermetallics and quasi-bina,:y Tb(Feo.aMo.2) 2 and Dy(Feo.sMo.2)2-based alloys with M = Co, AI, Si, Ga were subjected to the HDDR treatment. It was shown that at certain temperatures of HDDR treatment it is possible to increase substantially the coercive force of alloys, which allows the use of the obtained materials as cD,ogenic isotropic permanent magnets (for example, as bonded magnets). Such HDDR-magnets have the extended working interval of temperatures in comparison with those prepared from initial Dy(FeosMo.2)2 alloys. The HDDR treatment is shown to vitally influence the stepwise low-temperature magnetic reversal in Dy(Fe,AI)2. Due to the increase of both the coercive force and, consequent!y~ ~ heat release at lhe magnetization jump, it is impossible co obtain the isothermal conditions for the u~agr~tic reversal process for HDDP. Dy(Fe,Al)z at T < 6 K.

276

N.K. Zajkov, N.V Mushnikov I Journal of Alloys and Compounds 260 (1997) 271-276

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