Hydrogen-induced amorphization of intermetallics

Journal of

AHD COMPOUHDS ELSEVIER

Journal of Alloys and Compounds 231 (1995) 20-28

Hydrogen-induced amorphization of intermetallics K. Aoki*, T. Masumoto Institute for Materials Research, Tohoku University, Sendal 980-77, Japan

Abstract This paper describes our investigations on hydrogen-induced amorphization (HIA) in intermetallics caused by hydrogenation. The crystal structures and chemical compositions of amorphizing intermetallics are presented at first. In addition, the formation conditions and the formation processes of amorphous alloys, the mechanism of HIA, the factors controlling the occurrence of HIA and the relation between the structures of the hydrogen-induced amorphous alloys and the stability of C15 Laves phases are given. Keywords: Hydrogen; Amorphous; Intermetallic; Crystallization; Transformation; Thermodynamics; Magnetization

1. Introduction

2. Formation of amorphous alloys on hydrogenation

Hydrogen-induced amorphization (HIA), i.e. the transformation from the crystalline to the amorphous state caused by hydrogenation, was first demonstrated by Yeh et al. in the rapidly quenched metastable L12-type Zr3Rh in 1983 [1]. Subsequently, we reported HIA of the C15 Laves phases RM 2, (R = a rare earth metal, M = Fe, Ni, Co) [2,3]. The progress of HIA in the L12-type Zr3A1 compound was observed in situ by high-voltage transmission electron microscopy [4]. Although it had been suggested that amorphous alloys can be formed by hydrogenation of intermetallics [5,6], experimental results on HIA had been limited to the above alloy systems with the cubic (C15 and L12) structures. In order to understand the nature of HIA, it is useful to determine the types of amorphizing intermetallics and the formation conditions of amorphous alloys on hydrogenation. In this paper, the types (crystal structures, chemical compositions) of amorphizing intermetallics, the formation conditions and the formation processes of amorphous alloys are described. In addition, the mechanism of HIA and the factors controlling HIA in the C15 Laves phases and a correlation between the structure of the amorphous alloys and the stability of the C15 Laves phases are discussed.

2.1. Structural changes of intermetallic compounds on hydrogenation

*Corresponding author, 0925-8388/95/$09.50 © 1995 Elsevier Science S.A. All rights reserved SSDI 0925-8388(95)01832-8

As an example, powder X-ray diffraction (XRD) patterns of C15-type GdFe 2 hydrogenated at various temperatures and at 5 MPa H 2 for 86 ks are shown in Fig. 1 [7]. GdFe 2 absorbs hydrogen in the crystalline state at 300 K and the Bragg peaks shift to lower angles indicating the volume expansion of the lattice. In the XRD pattern of the sample hydrogenated at 423 K, however, the Bragg peaks disappear and are replaced by a broad maximum. The bright-field image of such a sample is featureless and the corresponding diffraction pattern shows a diffuse halo characteristic of amorphous alloys. Furthermore, a differential scanning calorimetry (DSC) curve shows an exothermic peak of crystallization when heated above 873 K. These observations indicate the amorphous nature of this sample. At higher hydrogenation temperatures above 673 K, new Bragg peaks corresponding to GdH 2 and a-Fe appear. Thus c-GdFe2H x (c indicates crystalline), a-GdFe2H x (a indicates amorphous) and GdH 2 and a-Fe are formed with increasing hydrogenation temperature. Fig. 2 shows thermomagnetization curves of a- and c-GdFe 2 with and without hydrogen prepared by several methods [8]. It is noticeable that there is a pronounced difference in the thermomagnetization

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

,

,

~

I

Hydrogenated at 5.0 MPa H2

i

oO '~ o •

o ~ O • ,~ A o/o

-

,

o ~

-

,~ j.~

~-Fe 7~ ^1 v

¢o

~,~ o ~ ~ 300 K , , . ~

Original GdFe2 ~ ,o~ ,~

~ .......•:_ ~ ~ 40 50 60 20/(zU180) tad

' 30

20

L.

423 K I ................ :::.=: :,.

~ ~ ~ :. g •~ ,~ ,o ~ ~,~L,,J~ ~ ~

L

[o GdH2

~ ool co ,~ o ~o4 ~ 673K o °

== ~E ~ ~

,

~ 70

80

Fig. 1. The powder X-ray diffraction patterns of GdFe 2 hydrogenated at various temperatures and at 5 MPa H 2 for 86 ks.

70 60

I~)'~ •-.,

(1; ~-G~. . . . o) c-O~,H,(373K~rCWa~) "~ (2) ~Gd.F~H~(473K,SMPa~) ",,. ~ o) ~O~°,(R03 .-" ..,.,_ ~ \ (~ ~C~F,~n,~cRW*7~UM~,n~)

50

'

'

~

(3) / : ~,~',, ~ \,/ / / ' - -,~.,,,, ...~ .-N 30 ;-'~. //"', \,,% /I'\----,,.\\ •2~i ./ ";,", / "~-.~-. 2~ ~'!~: ~,x,,/.~ "~..~ o 40

1~

%

/\

(~) ',

\.

/

,~'-~'b.

_ ~',"-3,

' - . ~ ..... J/\'--

100 2()0 300 400 500 6()()-"700 800 900 Temperature / K

2. The thermomagnetization curves of a- and c-GdFez with and without hydrogen prepared by various methods, Fig.

curves between a- and c-GdFe2H ~. The magnetization of c-GdFe2H4. 4 decreases drastically with increasing temperature around 100 K, but increases to the level of the original compound above 400 K. The recovery of the magnetization results from the desorption of hydrogen. The thermomagnetization curve of aGdFe2H3. 6 prepared by H I A is identical to that of

21

a-Gd33Fe67H120 prepared by hydrogenation of the rapidly quenched amorphous alloy. This result supports the similarity of the structure of both amorphous alloys. The difference in the thermomagnetization curves between a-GdF% and a-GdFezH x is small above about 300 K, suggesting that the magnetic properties of these amorphous alloys are dependent on their structures rather than hydrogen. Table 1 shows the crystallization temperature T x of the amorphous alloys, the Curie temperature T c as well as the amount of hydrogen absorbed and that desorbed after heating to 873 K in an argon atmosphere [8]. The same amount of hydrogen, 1.2 (H/M), expressed as the hydrogen-metal atom ratio, is absorbed in the rapidly quenched amorphous and the hydrogen-induced amorphous alloys, suggesting the similarity of their structures. Their hydrogen contents drop to 0.67 (H/M) releasing about a half of hydrogen, 0.53 (H/M). Hydrogen of 1.47 (H/M) is absorbed in the crystalline phase and becomes c-GdFe2H4. 4. It loses about two-thirds of hydrogen when heated to 873 K, (H/M) falls from 1.47 to 0.50 and 0.97 (H/M) of hydrogen is desorbed. Thus more hydrogen is absorbed and desorbed in the crystalline state of GdFe 2. These experimental results imply that hydrogen in a-GdFe2H x is trapped more tightly than that in cGdFe2H~, which may be closely related to the mechanism of HIA. Fig. 3 shows the thermomagnetization curves of c-CeFe 2 and a-Ce33Fe67 prepared by sputtering t0gether with the hydrogenated crystalline and the amorphous phases [9]. The magnetization of aCe33Fe67 is lower than that of c-CeFe 2. Hydrogen raises both the magnetization and the Curie temperature of the originally crystalline and amorphous alloys. The thermomagnetization curve of the hydrogenated crystalline sample is identical with that of the hydrogenated amorphous film within experimental accuracy, which also supports the amorphous nature of the hydrogenated crystalline sample.

2.2. Chemical compositions and crystal structures of amorphizing intermetallics Structural changes of A.B] x intermetallics (where A is a hydride-forming element such as rare earth metals, Zr, Ti, Hf, Ca and Mg, and B is a non-hydrideforming element such as ml, Ga, In, Mn, Fe, Co, Ni, Cu) caused by hydrogenation were examined by XRD, transmission electron microscopy (TEM), differential scanning calorimetry (DSC), the magnetic property measurement and hydrogen analysis. The chemical compositions and the crystal structures of the amorphizing intermetallics are listed in Table 2 [10-13]. As can bee seen in this table, more than 70 intermetallics can amorphize by hydrogenation. Amorphizing inter-

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

22

Table 1 The crystallization temperature Tx, the Curie temperature Tc as well as the amounts of hydrogen absorbed and those desorbed after heating to 873 K in an argon atmosphere Absorbed H 2

Alloys

(H/M) c-GdFeE(arc melted) c-GdFe2H4.4(373K, Hz) a-GdFeEHa.6(473K, HE) a-Gd33Fe67(RQ) a-Gda3Fe67HlEo(RQ + 473K, HE)

a¢

1.47 1.20 1.20

Tx (K)

Tc (K)

643 833 678

818 107 443 568 598

, '

~

--

~:---. "'%,", "°"o~.._ "~',~,~". ""o / N~'q a- CeFeEH.._ x -.°\j ,-.,.,. "* 'o \, ~.~4koe

oc 40

x~~

"~ g 2c

a-Ce33Fe+7

Fe"~X

0

, 10o

_ ~ -// "C-CcFc2 -~- ' -_" \ ~',~200

_ \,x# "x',.~-'

a-Cea~Fe67HxXA "

°--°-6 30o

Temperature / K

Z'~o , ~-,400

5oo

Desorbed H 2 (H/M) 0.97 0.53 0.53

In the AEB-type phases, C23 intermetallics such as REAl and B 8 E intermetallics such as REIn and ZrzA1 amorphize. In the ABE-type phases, only C15 Laves phases amorphize. Among the C15 Laves phases, RM 2 intermetallics (M = Fe, Co, Ni) amorphize, but RA12, ZrVE and others do not. The factors controlling the occurrence of HIA in the C15 Laves phases are discussed later. Thus HIA is confined to intermetallics with the specific crystal structures in contrast with the other solid-state amorphization reactions (SSARs). These experimental results suggest that HIA occurs when hydrogen atoms are in special environments, which is directly related to the mechanism of HIA.

Fig. 3. The thermomagnetization curves of a- and c-CeFe 2 with and without hydrogen prepared by various methods.

metallics have the A3B-, AEB- and ABE-type compositions, but no amorphization is observed in the AB-, AB 3- and A B : t y p e intermetallics. In the A3B-type phases, L12 intermetallics, such as Zr3M ( M = I n , AI, Rh) and R3In (R=rare earth metal), and D0~9 intermetallics such as R3Ga, RsAI and Ti3M (M = Ga and In) amorphize [14,15]. Ti3AI with the D019 structure does not amorphize, but (TiEZr)A1 and (TiEHf)Al with the same structure amorphize by hydrogenation,

3. Hydrogen-induced amorphization of the C15 Laves phase

3.1 HIA of RCo 2 phases investigated by differential thermal analysis (DTA) in a hydrogen atmosphere In order for HIA to occur, an amorphous phase must have a lower free energy than the corresponding crystalline one and a kinetic barrier must exist to prevent the formation of the equilibrium phase. We made thermal analysis of RCo 2 in a hydrogen atmos-

Table 2 The chemical compositions and the crystal structures of amorphizing compounds Composition

Crystal structures

Compounds

A3B

L12 (fcc)

Zr3In, ZrsA1, Zr3Rh R31n (R ~-Ce, Pr, Nd, Sm) RaGa (R ~-La, Pr, Nd, Sm) RaA1 (R---La, Ce, Pr, Nd) Ti3Ga, Ti31n (TiEZr)Al, (TiEHf)AI

D0t9

AEB

C23 B82

REAl (R-= Y, Pr, Nd, Sm, Gd, Tb, Dy, Ho) REin (R---La, Ce, Nd, Sm, Gd, Tb, Dy, Ho, Er) ZrEA1

AB 2

C15

RFe2 (R~Y, Ce, Sm, Gd, To, Dy, Ho, Er) RCo 2 (R---Y, Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er) RNi 2 (R---Y, La, Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er)

A, hydride-forming elements; B, hydride-non-forming elements.

K. Aoki, 7". Masumoto / Journal o f Alloys and Compounds 231 (1995) 2 0 - 2 8

phere to obtain information on the thermodynamic and kinetic aspects of i l i A . Fig. 4 shows D T A curves of C15 Laves R C o : (R = Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er) heated in a 1.0 MPa hydrogen atmosphere [16]. The D T A curves show three or four exothermic peaks. The X R D pattern of GdCo 2 heated above the first exothermic peak (to 420 K) exhibits the Bragg peaks which shift to l o w e r angles, indicating t h a t t h e f i r s t e x o t h e r m i c peak is due to the formation of c-GdCo2H X. In the X R D pattern of G d C o 2 heated above the second peak (to 610 K), the Bragg peaks disappear and are replaced by a broad maximum, indicating that the second exothermic peak is attributed to HIA. These D T A data clearly demonstrate that c - G d C o z H x is less stable than a-GdCoeH~, because the former transforms exothermically to the latter. The X R D pattern of GdCo 2 heated above the third peak (to 7 1 0 K ) is indexed on the basis of G d H 2 and/3-Co, indicating that the third peak is owing to crystallization. Similarly, the first, second and third exothermic peaks of RCo 2 are attributed to hydrogen absorption, H I A and crystallization, respectively.

,

,

RCo 2

,

,

1.0MPa H2 20K/min

ErGo 2

l

j

.~ ._ ,-, 2.

~ ~ ~ / ' ~ - - - ~

./,..../ D~,,~

/

-TbCO2

HoCo 2

0.7 0.6

0.5 ~ 0.4

RC02 ,,. ~.,,,.,.

23

1.0MPa H2 20K/min t,,,,',' , • , '

T#T,. _ PrCo2 __zx 2~2Z..SmCo2Gdco~N,O.-L^ _ T#T,. CeCo, - ~ . _ D ~ HeCo~ za " "'~;-A~_.Co~ PrCoz NdCo2 SmCo2GdCo2 . ......

e._o . . . . . . . . w •---~.92-DyCo,z-HoCo2--~~ • • •ErCo~'

v--'~ --_Pr,~oo k___"_~. 0.3 -Th/Tra . .N~Co2~,7~__~.GdC.o~SmCo2 .. mbCos _ _ ErGo.... o- - - o - - - _. ~~e~ ou ~__ * 0.2

[ O,

Hz absorp.

[

[ •

H~ Crystallization

I

O

0.1

z~

HoC°z

H2 absorp. & HIA

0 , , . a . , . , , , . i , , , i , , , I , , , 1100 1200 1300 1400 1500 1600 1700

Melting Point, T.,/K

Fig.

5. The ratio T , / T m vs.

T m for RCo 2 compounds.

Fig. 5 shows the T t / Z m ratios against the melting point T m of the original RCo e phases [16]. In the present work, the peak temperature was defined as the transformation temperature T,, i.e. the temperatures of hydrogen absorption Th, amorphization T. and crystallization T x. All these ratios exhibit a slight decreasing tendency with increasing Tm. However, it is worth noticing that T h / Tin, T a / T m and T x / T m are approximately constant, being 0.28, 0.40 and 0.50 respectively. These values are believed to be closely related to the mechanism of HIA, as will be discussed later. T

1

r

o ~

~

SmCo2

O LLI

1)

r DSC GdF%

H2 Absorp.

~

20 K/min

NdCo2 ~

PrCo2

HIA

~ ~

~

GdH~ + ct-Fe

~

Prec. of GdH,

., .t~p. / I 300

400

-

500 600 700 Temperature, T / K

800

Fig. 4. D T A curves of R C o z (R = Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er) heated in a hydrogen atmosphere at a pressure of 1 MPa.

3oo

400

500 600 Temperature TIK

/ 700

\" 800

Fig. 6. D S C curves of G d F e z heated in a hydrogen atmosphere of 1 MPa.

24

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

3.2. Correlation between stability of C15 Laves phases G d M e (11/1--- Fe, Co, Ni) and structures of aGdMeHx alloys

stability of the intermetallics and crystallization behaviors and structures of H I A alloys is discussed. Figs. 6 and 7 show DSC curves of GdFe 2 and GdNi 2 heated in a hydrogen atmosphere of 1.0 MPa respectively [17]. The DSC curve of GdFe 2 shows three sharp exothermic peaks and a broad one together with two endothermic peaks owing to hydrogen desorption. Similarly, the DSC curves of GdCo 2 show two sharp and two weak exothermic peaks together with two sharp endothermic ones (not shown). On the other hand, the DSC curve of GdNi~ shows three broad exothermic peaks together with a sharp endothermic one. The origin of the respective peaks, determined by XRD, T E M and hydrogen analysis, is indicated in Figs. 6 and 7. It is worth noticing that a-GdFe2H ~ (a-GdCo2H/) crystallizes to form an elemental hydride G d H 2 and pure metal Fe (Co), while a-GdNi2H ~ to G d H 2 and the intermetallic GdNi 5. From these data, the formation of the C15 Laves phase GdM 2 (M = Fe, Co, Ni) and their structural changes during heating in a hydrogen atmosphere can be summarized as follows.

As described in the previous section, D T A measurements showed that c-RCo~H x transforms to a-RCo2H x exothermically, which indicates qualitatively that the enthalpy of the former is higher than the latter. In this section, the enthalpy changes AH for hydrogen-related reactions, i.e. hydrogen absorption and desorption, H I A in the C15 Laves phase G d M 2 and crystallization of the resultant amorphous alloys are measured by DSC to understand the nature of HIA. Furthermore, structures of a-GdMeH x prepared by H I A are investigated by X-ray diffractometry. A correlation between

.

.

.

.

t e

Dsc GdNi~

t$

~

1.0 MPa H2 20 K/rain

H~ Abs0rp.

~-6 ~

/

~ ~

(I) For the G d - M - H 2 systems (M = Fe, Co)

A HIA

~ J

\¢

Gd + 2M +(x/2)H2--> c-GdMe +(x/2)H2--~ c-

St.

..............................

GdM2H ~ ~ a-GdM2H ~ ~ G d H 2 + 2M

H, Desorp.---~ I

300

I

400

(II) For the G d - N i - H 2 system

I

500

~0

700

Temperature T/K

s00

Gd + 2Ni +(x/2)H 2 --9 c-GdNi 2 +(x/2)H 2 ~ C-

Fig. 7. DSC curves of GdNi 2 heated in a hydrogen atmosphere of 1

G d N i 2 H ~ --->a - G d N i 2 H x ---~(3/5)GdH 2 + ( 2 /

MPa.

5)GdNi 5 + [(5x - 6)/10]H 2

Gd,2M,H2

Gd,2Fe, H2

Gd,2Co,H2

c-GdFe2 !:i: " ::::i

"ff -10{3 .:5 ~=

-14c

C-GdCo2 c-GdZe~HX

GdH2+2M

U -12t GdH2+2M** ,-i

Gd,2Ni,H:

GdH2+2Fe

......, ii!::::::!::FI2 Absorption

,

,i!i ' GdH2+2Co ' c-GdCChHx

. I . ~HC-GdNi2 c-GdNi2Hx

Formation of GdM2

-160 Amorphization

-180

a-GdNi2Hx

2+GdNi5

~Crystallization ,Ti

H2 Desorption

i

Fig. 8. Enthalpy changes for ~Lheformation of GdM 2, hydrogen absorption and desorption, HIA and crystallization along with the formation enthalpy of GdH 2.

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

In these equations, hydrogen desorption is described only in the last term, although in fact it occurs also in the transition from c-GdM2H x a n d / o r a-GdM2H x. The reactants are Gd, M and H 2 for every system. On the other hand, the products are GdH2, Fe(Co) and H 2 in G d - F e ( C o ) - H 2 systems, while they are GdH2, GdNi5 and H 2 in the G d - N i - H 2 system, Enthalpy changes for the total reaction is the same whether it takes place in one or several steps. Consequently, the total enthalpy change AH t for the reactions for G d - F e ( C o ) - H 2 systems should be equal to the formation enthalpy AHf of G d H 2. The values of AH for the formation of G d M 2, hydrogen absorption and desorption, H I A and crystallization are shown in the reaction sequence in Fig. 8 [17]. AHf for the formation of GdFe 2, GdCo 2 and GdNi 2 are - 35, - 48 a n d - 116 kJ mo1-1 Gd respectively [18]. The absolute value of AHf for the formation of GdNi 2 is about 2-3 times larger than that of the others, i.e. GdNi 2 is the most stable compound in the present C15 Laves phases. The sum of AH in the G d - F e ( C o ) - H 2 system is in good agreement with AHf of G d H 2. However, the sum of AH in the G d - N i - H z system shows a large negative value ( - 173 kJ m o l - l Gd). This result is not unusual, because a-GdNi2H x does not crystallize into G d H 2 and Ni. AHf of GdNi 2 is more negative than AHf of G d H 2. In addition, hydrogen absorption, H I A and crystallization occur exothermically. Therefore, the sum of AH are always more negative than AHf of G d H 2 in the G d - N i - H 2 system. Consequently, aGdNi2H . never crystallizes into G d H 2 and Ni. The sum of AHf and AH h in the G d - C o - H 2 system is more negative than AHf of G d H 2. However, hydrogen desorption can occur endothermically so as to raise AH. Therefore, it is possible for a - G d C o z H . to crystallize to form G d H 2 and Co. The present work indicates that the amorphous alloys prepared by hydrogenation of the less stable intermetallics such as c-GdFe 2 and c-GdCo 2 crystallize into the elemental hydride G d H 2 and pure metal Fe(Co), but that prepared from the stable compound GdNi 2 crystallizes to G d H 2 and GdNi 5. Therefore, we can conclude that the crystallization products of a-GdM2H x are determined by the stability of c-GdM 2. Next, we discuss a correlation between the stability of c-GdM 2 and the structure of a-GdM2H ~. The radial distribution functions (RDFs) of a-GdM2H x are shown in Fig. 9 [17]. The coordination number, N, and the nearest neighbor distance, r, calculated from these curves, are tabulated in Table 3, together with the values calculated from a - G d F e 2, c-GeFe 2, G d H 2 and the Goldschmidt diameters [17]. r (0.381 nm) of the G d - G d pair for a-GdFeEH 3 is larger than that (0.347 nm) for a-GdFe2, which suggests that hydrogen atoms occupy the sites surrounded by Gd atoms interstitially in a-GdFeEH 3. N for a - G d F e 2 indicates that Fe atoms

25

800 600 --u_ rr v ot'"-= e"~ t.t_ e.2 ~ ~ .E a -~ ~ tr

400

~ ~_ ~ +

200 GdFe2H3.0

0 ~ o oo ~ ~ o

400

200 0

o1

r,,.

/~/

GdCo2H3.4

"/-

400 200

. GdNi2H3

._ / z ~ ~ ~ J ~1 1/~/---~"

/

y -

0 _2000

i .1

i .2

i .3 r

i . ~,

i .5

I .6

.7

/nrn

Fig. 9. The radial distribution functions (RDFs) of a-GdMEH x alloys prepared by hydrogenation.

have an average of 6.3 Fe and 3.3 G d nearest neighbors; G d atoms have an average 6 G d nearest neighbors. Correspondingly, N values in a-GdFe2H3. 0 are 6.7, 1.3 and 11.4. These indicate that there exists a greater tendency for the presence of F e - F e and G d Gd nearest neighbor pairs in a-GdFe2H 3 than in aG d F e 2. In particular, N = 11.4 of the G d - G d pair for a-GdFe2H 3 is almost the same as that for c - G d H 2, which suggests that the short-range order in aG d F e 2 H 3 is similar to that in c - G d H 2. N values of the M - M and G d - G d pair increase in the order of aGdNi2H3. 5, a-GdCo2H3. 4 and a-GdFe2H 3. On the contrary, N values of the M - G d pairs increase in the reverse order. Both larger N values of the M - M and G d - G d pairs and small N values of the M - G d pair indicate the strong tendency for M and Gd atoms to d u s t e r in a-GdFe2H 3 and a-GdCo2H3. 4. As mentioned before, a - G d F % H , and a - G d C o 2 H ~ crystallize to form G d H 2 and Fe(Co). Such crystallization behavior is considered t o correspond to the clustering of Fe(Co) and G d atoms. N of the M - G d pair for aGdNiEH3.5, a-GdCoEH3. 4 and a-GdFeEH ~ indicate that Ni, Co and Fe atoms have an average of 4.3, 2.4 and 1.3 G d nearest neighbor, respectively. This implies that the pair interaction energy ENi_GO between Gd

26

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

Table 3 The coordination number N, and interatomic distance r, in a-GdM2Hx alloys prepared by hydrogenation together with the values calculated from a-GdFez, c-GdFe2, GdH 2 and the Goldschmidt diameters Alloys a-GdFe2H3.0 a-GdFe 2 c-GdFe2 Goldschmidt diameter a-GdCoEH34 a-GdNi2H35 c-GdH2

M-M

M-Gd

Gd-Gd

r(nm)

N

r(nm)

N

r(nm)

N

0.258 0.254 0.260 0.254 0.256 0.259

6.7 6.3 _+0.5 6

0.313 0.307 0.305 0.307 0.310 0.311

1.3 3.3 _+0.3 6

0.381 0.347 0.318 0.360 0.375 0.387 0.375

11.4 6 _+1 4

6.2 5.2

and Ni atoms is the largest, c-GdNi 2 is the most stable compound in the present C15 Laves phases. Consequently, it is conside, red that the strong correlation between M and Gd atoms is present in the amorphous alloys prepared by H I A of the stable intermetallic, The present work concludes that the structure of the amorphous alloys prepared by hydrogenation of the stable GdNi 2 is homogeneous rather than that prepared from the less s'Lable ones. 3.3. Factors controlling the occurrence o f H I A in the C15 Laves phases

Not all compounds with the crystal structures shown in Table 2 amorphize on hydrogenation. For instance, the C15 Laves phase RM 2 (R = a rare earth metal, M = Fe, Co, Ni) amorphize by hydrogenation around 400-500 K, while ZrV2, Z r M o 2 and others absorb hydrogen retaining the crystalline state. Consequently, H I A is controlled not only by the crystal structure, but also by other factors. In the present section, the factors controlling the formation of the amorphous alloys will be analyzed in terms of the atomic size ratio, the electron concentration, the thermal stability of original compounds, the size of tetrahedral holes occupied by hydrogen atoms, etc. The AB 2 Laves phases are so-called size factor compounds and are constructed from the close packing of the hard spheres with the ideal atomic size ratio R A / R B of 1.225. Substantial mutual adjustments of the atomic size can take place when the compounds are formed by atomic species whose radius ratio does not coincide with this ideal value. The parameters (rA RA) and (r B - R B ) are employed as indices of the size adjustments. Here, r A and r B are the radii of A and B atoms calculated from the lattice parameters and the packinj~ geometry of the C15 Laves compounds, i.e. r A = x/3a/8, r B = x/'2a/8, and R A and R B are the Goldschmidt radii of A and B atoms. Structural correlations of hydrogenated C15 Laves phase compounds are examined by plotting the parameters (r A -- RA) and (r~ - RB) against R A / R B in Figs. 10 and 11 respectively [7]. The values of (rg - R A )

2.4 4.3

7.4 4.9 12

decrease with increasing R A / R B ratio, i.e. the A atoms contract as the atomic size ratio becomes large. In contrast, the values of (r B - RB) increase with increasing R A / R B values. The B atoms exhibit size contraction ((r a - - R a ) < 0) when R g / R a is smaller than 1.37, but show an increase in effective size ((r a - - R B ) > 0) when R A / R a is larger than 1.37. H I A occurs under such conditions. Thus both the large contraction of the A atoms and the expansion of the B atoms are closely related to the occurrence of HIA. It is concluded that the atomic size ratio is the most important factor controlling the occurrence of H I A in the C15 Laves phase AB 2 and the intermetallics with the ratio R g / R a larger than 1.37 are amorphized by hydrogenation. That is, strained C15 Laves phases are apt to amorphize by hydrogenation. The well-known empirical rule for the size effect states that the atomic radii of two elements must differ by more than 10% in order for their binary alloys to form amorphous alloys by rapid quenching. The difference in the atomic sizes (37%) required for the occurrence of H I A is much larger than that required for the formation of the amorphous alloys by rapid quenching. 3.4. The m e c h a n i s m o f H I A in the C 1 5 L a v e s p h a s e s

Finally, we discuss the thermodynamic driving force and the kinetic aspect of H I A in the C15 Laves phases. The D T A and DSC data demonstrate that the enthalpy of a-RM2H x is lower than that of c-RM2H x, because the latter transforms to the former exothermically. The driving force for H I A in c-RM 2 is considered to result from the enthalpy difference in the two states of the alloys. The enthalpy difference is explained on the basis of the environmental difference of hydrogen atoms in a- and c-RM2H x as follows [16]. Hydrogen atoms in the C15 Laves c-RM2H x occupy the tetrahedral sites surrounded by 2R + 2M and 1R + 3M according to the geometrical constraints [19]. On the contrary, hydrogen atoms in the corresponding a-RM2H x occupy the tetrahedral sites surrounded by 4R and 3R + 1M and others. Since the formation enthalpies of R H 2 have lager negative values, hydro-

K. Aoki, T. Masumoto I Journal of Alloys and Compounds 231 (1995) 20-28

27

C15 Laves phases

z,v2 o

0

N <

I, I

\riCo2 o ~

1.225

o

2

O

fCo2

-0.010

"~

I I

=

I

o

~-~ I

SmAl20:3~Co2

E

"~= -

-Ea2~2

~,~

CeNi,'2.

I

"-g

]

•

, ~

o~.o~

kaPt2. •~. "t-ioFe2 O ~)v~',;~• ~ S m F e 2

-0.020

, ~.,.~ ,c~. a~ .o~

12

.~

[i Amorphous

'~

=v"~vr~

~ Crystalline

I

I ,

I

,

I

,

~

I.I0

I

,

,

,

,

120

l

,

,

~SmCe2

GdN'e•'~rCO2prN~------------------~x

I -0.030

,=

t,J~tSO z.~ ~ , N - " b C d 2

\

1.37

, I ,

150

1

,

,

,

,

1.40

Atomic Size Ratio

e LaN'~. l

150

RA/R B

Fig. 10. The structures of the hydrogenated C15 Laves phases against (rA -R~) and Ra/R~.

C15 Laves phases '

I

'

'

'

'

I

. . . .

I

'

'

'

i

I

'

I

. . . .

I

'

I

I iI

;Amorphous

Crystalline

I

O.OLO

GdMn2

l 1

II

Dy~ Lag~

I I

E

1.225

I

0

I

=

•I

'

~

•~ "~"

,..,PrFe.-~ PL-CO2

C:,AIglLaAI2

eFe~CeC°~

Laa~ ~

ZrCo2 o 1 ~

'

_.

ZrFe20?.,2gO ~ t ' t 2 0

--~

]

I

/

O~Co2

O..,

ScAt2 O

E~ o

43.010

o

1.37 ,

I

1.10

,

,

,

,

I

,

,

,

120

,

I

,

,

,_

I_30

Atomic Size Ratio

,

I

1.40

,

,

,

,

I

,

1.50

RA/R B

Fig. 11. The structures of the hydrogenated C15 Laves phases against (rB -RB) and RA/R B.

gen a t o m s in the a m o r p h o u s phases are m u c h m o r e strongly b o u n d t h a n those in the crystalline phases. In o t h e r words, h y d r o g e n atoms in the crystalline phases are less energetically favorable. F o r instance, if c-

R C o 2 H x are h e a t e d to the t e m p e r a t u r e s (0.4Tin) as shown in Fig. 5 w h e r e the metallic atoms can m o v e over a short distance, the r e a r r a n g e m e n t of the metallic a t o m s can occur to r e d u c e the e n t h a l p y of c-

28

K. Aoki, T. Masumoto / Journal of Alloys and Compounds 231 (1995) 20-28

RCo2H x, which leads to HIA. At the lower temperaturPs (0.28Tm) as shown in Fig. 5 where the diffusion rates of the metallic atoms are very low, c-RCo 2 absorbs hydrogen in the crystalline state. That is, H I A is suppressed for kinetic reasons at lower temperatures. On the contrary, at higher temperatures (0.5Tm), where the metallic atoms can move over a long distance, the amorphous phase is no longer stable and decomposes into R H 2 and Co. Thus, crystallization occurs above 0 . 5 T m.

4. Summary and conclusions Structural changes of AxBI_ x compounds (where A is a rare earth metal, Zr, Ti, Hf, Ca or Mg, and B is AI, Ga, In, Mn, Fe, Co, Ni, Cu) on hydrogenation were examined by XRD, TEM, DSC, the magnetic property m e a s u r e m e n t a n d hydrogen analysis. The amorphizing intermetallics are L12, D019 , C23, B82 and C15 phases with AB2-, A2B- and A3B-type compositions. Hydrogenated crystalline phases c-RM2H x transform t o a RM2H x exothermally. As the melting point T m of the C15 Laves phase RCo 2 increases, the hydrogen absorption temperature Th, H I A temperature T a and crystallization temperature T x increase. On the contrary, T h / T m , T a / T m and T x / T m a r e nearly constant, being 0.28, 0.4 and 0.5 respectively. This result suggests that H I A occurs by the diffusion of metallic a t o m s o v e r a short distance. The amorphous alloys prepared by hydrogenation from the less-stable C15 Laves GdM 2 show the strong tendency for the s a m e kind of metal a t o m s t o cluster and crystallize into elemental hydride G d H 2 and the pure metal M. Correspondingly, the sum of the enthalpy change of the total reaction becomes equal to that of the formation enthalpy for GdH 2. On the contrary, the a m o r phous phase prepared from the stable intermetallics shows a rather homogeneous structure and crystallizes into GdH 2 and an intermetallic. The atomic size ratio is the decisive factor controlling the occurrence of H I A in the C15 Lawns AB 2 phase, and those with

ratios over 1.37 amorphize by hydrogenation. The thermodynamic driving force for H I A in the C15 Laves compounds is considered to be the enthalpy difference resulting from the different hydrogen occupation sites in the two states of the alloys.

References [1] X.L. Yeh, K. Samwer and W.L. Johnson, Appl. Phys. Lett., 42 (1983) 242. [2] K. Aoki, K. Shirakawa and T. Masumoto, Sci. Rep. Res. Inst. Tohoku University, A-32 (1985) 239. [3] K. Aoki, T. Yamamoto and T. Masumoto, Sci. Rep. Res. Inst. Tohoku University, A-33 (1986) 163. [4] H. Fujita (ed), Proc. Int. Syrup. on Behavior of Lattice lmperfections in Metals--in situ Experimentation with HVEM, Osaka, Japan, 1985, p. 1. [5] H. Oesterreicher, J. Clinton and H. Bittner, Mater. Res. Bull., 11 (1976) 124. [6] I. Jacob and D. Shaltiel, J. Less-common Met., 65 (1979) 117. [7] K. Aoki, X-G. Li and T. Masumoto, Acta MetalL Mater., 40 (1992) 1717. [8] K. Aoki, M. Nagano, A. Yanagitani and T. Masumoto, J. Appl. Phys., 62 (1987)3314. [9] K. Aoki, T. Yamamoto, Y. Satoh, K. Fukamichi and T. Masumoto, Acta MetaU., 35 (1987) 2465. [10] K. Aoki, T. Yamamoto and T. Masumoto, Scr. Metall., 21 (1987) 27. [11] K. Aoki, A. Yanagitani, X-G. Li and T. Masumoto, Mater. Sci. Eng., 97 (1988)35. [12] K. Aoki, X.-G. Li, T. Aihara and T. Masumoto, Mater. Sci. Eng., A133 (1991)316. [13] K. Aoki and T. Masumoto, J. Alloys Comp., 194 (1993) 251. [14] K. Mori, K. Aoki and T. Masumoto, Scr. Metall. Mater., 27 (1992) 623. [15] K. Mori, K. Aoki and T. Masumoto, Sci. Rep. Res. Inst. Tohoku University, A-38 (1993)43. [16] K. Mori, K. Aoki and T. Masumoto, Mater. Sci. Eng., A 179/180 (1994) 181. [17] K. Aoki, X.-G. Li T. Hirata, E. Matsubara, Y. Waseda and T. Masumoto, Acta Metall. Mater., 41 (1993) 1523. [18] C. Colinet and A. Pasturel, J. Less-Common Met., 119 (1986) 167. [19] D. Ivey and D. Northwood, J. Less-Common Met., 115 (1986)

23.