Metal hydrides: A review of group V transition metals—niobium, vanadium and tantalum

Int. J. I~vdrogenEnergy, Vol. 17, No. I, pp. 41-52, 1992.

0360-3199/92 $5.00 + 0.00 Pergamon Press plc. © 1992 InternationalAssociation for Hydrogen Energy.

Printed in Great Britain.

METAL HYDRIDES: A REVIEW OF GROUP V TRANSITION METALS--NIOBIUM, VANADIUM A N D TANTALUM A. Y. ESAYED and D. O. NORTHWOOD Engineering Materials Group, Mechanical Engineering Department, University of Windsor, Windsor, Ontario, Canada N9B 3PB

(Received for publication 5 November 1991) Abstract--The group V transition metal (Nb, V, Ta)-hydrogen systems are reviewed. Consideration is given to the different types of hydrides formed, their crystallography, thermodynamic data and hysteresis effects on hydriding/dehydriding. The group V transition metals do not themselves react with hydrogen at room temperature, but quite readily react with hydrogen when alloyed with other bcc metals of smaller atomic radii. There are uncertainties regarding the structures of hydrides formed by this group of metals. These uncertainties mainly arise because of the limited number of neutron diffraction studies of the hydrides. Monohydrides have been reported for Nb, V and Ta, while dihydrides have been reported only for Nb and V. Thermodynamic data show that the niobium-hydrogen solid solution is more stable than the vanadium - or tantalum-hydrogen solid solutions. Hysteresis have been seen in metal-H systems in this group and, as yet there is no complete explanation for its occurrence.

2.1. Monohydrides

INTRODUCTION

The monohydrides of group V transition metals can be formed by heating the metal at moderate temperatures ( 3 0 0 - 7 0 0 ° C ) and pressures in a hydrogen atmosphere. Albrecht and co-workers [4] studied the rate of reaction of hydrogen with niobium over the temperature range 3 0 0 - 7 0 0 ° C . They found that between 600°C and 700°C the reaction followed a parabolic rate law, i.e the reaction is diffusion rate controlled. In addition, they found that at temperatures between 300 and 500°C the surface saturation necessary for a parabolic rate behaviour was not obtained. They concluded that this behaviour was due to some unexplained surface reaction. Mallett and Koehl [5] in their investigation of the rate of reaction of hydrogen with tantalum at temperatures of 500°C and above, also found that the rate was parabolic. The monohydride can also be made by heating a mixture of Call2 and either Nb205 or Ta205 to temperatures not exceeding 900°C, and then slow cooling down to room temperature. The hydride is formed during the cooling process [ 6 ] . Formm and Uchida [2] have also found that the monohydrides of group V transition metals can be formed by heating the metal to high temperatures ( 1 3 0 0 - 2 2 0 0 ° C ) in an ultra-high vacuum ( < 10 -9 Pa) followed by exposure to hydrogen at the same temperatures.

The group V metals Nb, V, and Ta (transition metals) which have the body-centred cubic (bcc) structure do not readily form hydrides. In bulk form, the reaction will not occur at room temperature. This is rather surprising given the fact that the diffusion of hydrogen in bcc metals is orders of magnitude higher than in other metals [1]. Formm and Uchida [2] have attributed the non-reaction at room temperature to an oxide layer which forms on the surface of the metal and both prohibits the catalytic dissociation of H2 and blocks H entry into the metal. Athough the group V transition metals do not readily form hydrides at room temperature, they react easily with hydrogen when alloyed with other bcc structure metals having smaller atomic radii. For instance, it has been reported that [ 1, 3] alloying Nb with V, Cr, Fe, and AI induced Nb to react rapidly with hydrogen at room temperature. However, alloying Nb with metals having equal or larger atomic radii than Nb such as Ta, Zr, or Ti had no effect on the reaction of Nb with H2. This paper is intended to serve as an overview of the different types of hydrides formed, their crystallography, various thermodynamic properties and hysteresis in m e t a l - h y d r o g e n systems of the group V transition metals (niobium, vanadium and tantalum).

2,2. Higher hydrides and dihydrides The dihydrides were first made by Brauer and Mueller [ 7 ] . They treated niobium metal powder with 10% hydrofluoric acid, and observed another phase, i.e. the dihydride, besides the monohydride. The same authors [8]

2. TYPES OF HYDRIDES This section is divided into two subsections, one dealing with monohydrides, and the other dealing with the higher hydrides, particularly the dihydrides. 41

42

A.Y. ESAYED and D. O. NORTHWOOD

prepared the dihydride by cathodically charging the niobium metal. The homogeneity range of the dihydride extended from NbH2.00 to NbH2.07. However, they reported that the niobium dihydride was unstable and decomposed into/3-NbH and hydrogen. Niobium dihydride has also been prepared by treatment of Nb205, or preferably NbH, with 10% hydrofluoric acid [9]. Maeland and co-workers [ 10] reported the formation of a non-stoichiometric dihydride of vanadium by crushing vanadium monohydride under high hydrogen pressures, 70 atm. The hydride they obtained had a composition of V H ~ . Examination by X-ray powder diffraction techniques indicated the presence of an fcc phase along with the body centred tetragonal (bct) monohydride. They calculated the lattice parameter of the fcc phase and found it to be 4.270 +0.002 A. They believed that the hydrogen atoms occupied the tetrahedral sites in the fcc vanadium lattice thus giving the new compound a fluorite type structure at VH2. The same authors found that the hydrofluoric acid treatment yielded a vanadium hydride of composition VH177. Brauer and Mueller [8], however, reported that the hydrofluoric acid treatment of Ta205 produced only Tall09. Samsonov and Antonova [ 11 ] prepared a higher hydride of niobium, Nb~.34, by exposing the niobium metal powder to hydrogen gas at 600°C. The same authors concluded that this was a metastable phase since this hydride transformed to a monohydride at temperatures above 1000°C. 3. CRYSTALLOGRAPHY

Brauer and Hermann [9] used X-ray diffraction method to study the hydrides and deuterides of niobium formed at room temperature. They found that the bcc structure exists up to compositions of about NbH(D)0.~. The two-phase field, i.e. Nb + hydride, extended to NbH(D)07. They also found that at higher hydrogen contents (H/Nb > 0.8) there was only the single phase hydride (deuteride) with a slightly distorted body-centred orthorhombic (bco) structure. structure can be indexed as a face-centred orthorhombic (fcc) structure with ao = 4.84 ~,, bo = 4.90 A and co = 3 . 4 5

/k.

Paxton et al. [ 15] examined niobium single crystals containing hydrogen that had been cooled through the phasetransformation temperature. The hydrogen was introduced at 1 atm and at temperatures of 300, 600 and 900°C. The samples were then cooled down to room temperature. They concluded that the transformation from the single phase to the two-phase condition was of the martensitic type. Wainwright [16] also used X-ray diffraction techniques to investigate niobium hydride with compositions between HbH0 t (hydrogen in solid solution) and NbH09 (hydrides) at various temperatures. Heating the orthorhombic, NbH0.9, hydride caused the distortion to decrease with increasing temperature such that at temperatures above 100°C the hydride was bcc. Moreover, when the temperature was increased, the lattice parameter of each phase in the two-phase region was dependent on the percentage of each phase that was present. This violates normal thermodynamic criteria and has led to the suggestion that a martensitic type of transformation is present in this system [6]. Wainwright et al. [17] using X-ray line

This section is divided into four sub-sections, the first three dealing with the hydrides of Nb, V and Ta respectively. The last sub-section is a summary comparing the hydrides formed by the three different metals.

lOO 100 °C

3.1. Niobium (columbium) hydride

Niobium metal has a bcc structure with a lattice parameter ao = 3.3066 A [6]. The change in lattice parameter as a function of hydrogen content and temperature was first determined by Albrecht and coworkers [4] who found an increase of about 0.0023 -4- 0.0002 ,~, in the lattice parameter per at% of hydrogen. They identified the hydride phase by X-ray diffraction measurements to be a bcc structure. Birnbaum and co-workers [13] studied the morphology of the monohydride H-phase, NbH075, in niobium and they found that both the cooling rate and degree of under-cooling effected its morphology. Rapid cooling caused a transition to a massive form of the hydride whilst precipition for a large degree of undercooling resulted in a "dendritic" morphology. Komjathy [13] investigated the hydride or, H-phase, of niobium by using X-rays and found the metal sublattice in this phase was bcc with a lattice parameter which varied with hydrogen content from 3.42 to 3.445 ~,. This is in a good agreement with the results of Albrecht et al. [4]. Reilly and Wiswall [ 14] ~howed that the niobium forms a dihydride (-y-phase) having the cubic fluorite structure.

10

/P/-

54 °C

g N .~

1.o

~5

I

Atomic-ratio H / V

Fig. 1. Pressure-composition isotherms for the VHI 0-VH2.0 [141.

GROUP V METAL HYDRIDES profile analysis showed that the hydride phase was a distorted cubic type at all compositions including those in the two phase region. 3.2. Vanadium hydrides Vanadium metal has a bcc structure with a lattice parameter ao = 3.04 ,~, [6]. Vanadium metal is almost inert to hydrogen at both room temperature and higher temperatures [ 18]. Reilly and Wiswall [14] found that vanadium forms dihydrides (-r-phase) having the cubic fluorite structure. Figure 1 shows a number of their pressure-composition isotherms for the region extending from VH-0s to VH-20. The initiation of a pressure plateau at a composition aproximating VH-095 marks the point at which -r-phase appears with/3-phase. At 40°C the plateau extends from VH0.95 to VH20 at which point the isotherm rises sharply marking the composition at which the ~-phase disappears and only ",/-phase exists. Lynch and co-workers [ 19] reported that the vanadium, as well the other group V transition metals, forms non-stoichiometric monohydrides (~-phase) which have distorted bcc structures. The preparation of a non-stoichiometric dihydride of vanadium has been reported [ 10, 20]. The dihydride had a fcc arrangement of vanadium atoms with a,, =4.270 ± 0.002 A . The hydrogen atoms presumably occupied tetrahedral sites resulting in fluorite type structure at VH2. This structure has not been demonstrated, but rather inferred from the estimated radius of hydrogen. Trzeciak et al. [21] reported that the ~-phase hydride, VH094, in the vanadium-h~'drogen system is bet with a. = 3.02 ,~ and co = 3.36 A . Roberts [22] used a neutron diffraction technique to determine the structure of WD0.75. The VD0.75 had a bcc structure for the metal atoms with ao = 3.138 ~,. The neutron-diffraction data were difficult to interpret. At higher temperatures, i.e. above 207 K, the deuterium atoms did not appear to be ordered, while at lower temperatures, below 207 K, they appeared to arrange themselves into a primitive cubic lattice of a,, = 6.30 A , which is about double the lattice parameter of the vanadium sublattice. 3.3. Tantalum hydrides Tantalum metal has a bcc structure with ao = 3.303 [6]. Tantalum does not readily absorb hydrogen at room temperature. However, tantalum readily takes up hydrogen at elevated temperatures ( > 300°C) to form a hydride with cubic symmetry [23]. Reilly and Wiswall [ 14] reported that the Ta does not form dihydrides (3,-phase) with the cubic fluorite structure like V and Nb. They found most of tantalum hydrides are non-stoichiometric monohydrides (~-phase) which have distorted bcc structures. Brauer and Mueller [8] reported that the hydrofluoric acid treatment of Ta205 gave only Tall09. We would suggest that the high pressure method, or the hydrofluoric acid treatment of Tall, might perhaps produce a higher hydride of tantalum. A neutron diffraction investigation by Wallace [24] has

43

shown that the hydrogen in Ta resides in the tetrahedral interstices of which there are six per unit cell. The theoretical composition is therefore Tall6 if all sites are filled. In Ta2H, which has been extensively studied only 1/12 of the available sites are filled. 3.4. Summary X-ray diffraction studies show the hydrides of vanadium and tantalum have a bet structure at low hydrogen contents close to the two-phase (o~ + ~) region. This can be viewed as a light distortion of the bcc metal lattice (Fig. 2a) which explains the continuous transition from metal to hydride at low temperatures. However, for the case of the vanadium system, both a solid solution phase and the ~-phase, which is a tetragonal hydride phase, have been identified. In the case of tantalum, the hydride becomes orthorhombic (Fig. 2b) as the hydrogen content increases. In the case of niobium the X-ray investigations show that the ~-phase occurs along with a phase which is orthorhombic at room temperature and has a composition of NbH0.8-09. At compositions below NbH0.8-0.9, the hydride remains cubic. A summary of the lattice parameters and structures of the hydride phases for all three of the group V transition metals is given in Table 1. There is still considerable uncertainty as to the structures of the monohydrides of the group V metals. This uncertainty is principally due to the limited number of neutron diffraction studies carried out on this group, as well as with difficulties in interpreting the available data. These interpretation difficulties arise because the monohydrides of this

a

Fig. 2. Structures of hydrides of group V transition metals: (a) distorted cubic structure a = c; (b) orthorhombic cell.

w

w

Fig. 3. Fluorite type metal dihydride structure. The filled circles represent metal atoms, the dashed open circles represent hydrogen atoms.

A. Y. ESAYED and D. O. NORTHWOOD

44

Table 1. Structures and lattice parameters of hydrides of the group V transition metals Nb, V and Ta

Metal Nb

v

Metal structure

Lattice parameter A

Hydrides

Hydrides structure

bcc

3.3066

NbH

o~horhombic

NbH2

~c

VH

bct

VH2

~c

TaH

bct

bcc

Ta

dcc

3.04

3.303

o~horhombic

group are bcc (to a first approximation) and as such there are six tetrahedral sites and three octahedral sites per unit cell. Therefore, the possibilities of disorder and superstructure formation are high. This compares to other groups in the periodic table, such as VI and VII, where the monohydrides are close packed with only two tetrahedral sites and one octahedral site per unit cell. The dihydrides of niobium and vanadium have a fcc structure (fluorite type, see Fig. 3). No dihydride has been reported for tantalum.

M + Y/2H2 = MHr.

(1)

The transition to M H r is the absorption process, and in general is an exothermic reaction, the heat of formation approaching the heat of combustion of hydrogen compounds. The reverse reaction is the desorption process. The behaviour of m e t a l - h y d r o g e n systems is most conveniently represented by a means of p r e s s u r e - c o m p o s i t i o n - t e m p e r a t u r e diagram ( P - C - T ) . The diagram is a plot of hydrogen pressure vs composition at various temperatures as shown in Fig. 4. At low hydrogen concentrations there is a strong composition dependence of hydrogen pressure. This region (i.e. the a-phase) corresponds to hydrogen going into solid solution without the occurrence of a second phase. When the curve starts to change slope on the P - C - T diagram this indicates the start of next 'stage' which involves a region where the pressure is independent of the hydrogen concentration. In

4.840 4.900 3.450 4.563

[10]

a = 3.020 c = 3.360 a = 4.270

[24]

a c a b c

= = = =

= = = = =

3.380 3.410 4.730 4.780 3.430

[9]

[11] [28] [10]

03

o)

o

E

g

The reaction of hydrogen with a metal M to form a stable hydride can be described by the following reaction (i.e. a direct reaction of metal with hydrogen gas):

a b c a

Ref.

(1)

~3

4. T H E R M O D Y N A M I C PROPERTIES

Hydrides lattice parameter A

LIJ

/

Absorption Desorption

~+13

I I I I

Hydrogen to metal ratio Fig. 4. Representative pressure-composition isotherm [30].

this region the saturated solid solution of hydrogen in the co-phase is in equilibrium with a hydrogen-deficient hydride phase, the 15-phase. As the second phase (/3-phase) forms, the pressure remains constant and a "plateau" results, as more hydrogen is added. Ftuther increases in hydrogen concentration after the ct-phase has been completely converted into the ~-phase give rise to an increase in pressure. An additional hydride phase, the -},-phase, may also be formed, in which case a second and higher plateau will appear. The standard enthalpy of formation, AHf, of a metal hydride M H r can be calculated from van't H e f t equation, (~i In Kp)/dT = A H f / R T 2

(2)

GROUP V METAL HYDRIDES

45

2

where R = gas constant T = temperature (K)

1

Kp = the equilibrium constant and is given by Kp = aM.2 /aMY(,, 12

(3) 0

or

Kp = (aMH2)/aMxa~ 2

(4) [3..

which becomes

_.qo Kp = p~V/2

(5)

when the standard states of hydrogen and metal are taken as the pure solids in each case (aM.2 = aM = 1) and all2 = PH2. Substituting equation (5) into equation (2) and then integrating (assuming AHf to be constant over a relative small temperature range), the van't Hoff isobar can be written as:

-2

..3

i

i

i

i

9

10

11

12

(10 4) / T(K)

d ln(P.2 v/:) = ( A H f / R T ) d T

(6)

Fig. 5. Logarithm of dissociation pressure vs I/T (K) for lutetium dihydride [30].

In P.2 = ( 2 / x ) ( A H / R T ) + C

(7)

Therefore, these metals do not readily react with hydrogen at room temperature, especially in bulk form [1, 6].

or

where C is the constant of integration. The enthalpy of formation of the hydride is calculated from the slope of the straight line made by plotting In P.2 vs 1/T (Fig. 5). The standard free energy of formation ,SG~' of the metal hydride can be determined by the relation: AG~' = - R T In Kp = (Y/2) R T In PH:.

(8)

The standard entropy of formation can then be determined by: 2xS~' = ( AH~' -- ,SG~')/T

(9)

where AHtY' = standard enthalpy of formation [25, 26]. For most metal-hydrogen systems where there is appreciable non-stoichiometry, the standard enthalpy of formation is the sum of three components, namely: (i) the integral heat of solution of hydrogen in the c~-phase from zero hydrogen content to saturation, (ii) the heat of reaction in going from the hydrogen saturated a-phase to the nonstoichiometric f3-phase, and (iii) the integral heat of solution of hydrogen in the hydrogen-poor-/3-phase up to the stoichiometric value. In the case where there are large deviations from stoichiometry, the thermodynamic quantities are usually expressed as relative molal quantities (Sf. - 1/2Xr~2 ), where Xf is the partial molal enthalpy (entropy or free energy of hydrogen as atoms in the solid) and X~2 refers to hydrogen in its standard state as a pure, diatomic, ideal gas [25]. For the group V transition metals, the free energy of reaction (1) at room temperature is strongly negative.

4.1. N i o b i u m - h y d r o g e n system

The terminal solubility of hydrogen, i.e. the maximum solubility of hydrogen in the metal phase (a-phase) before formation of a hydride, in Nb and Ta alloys was studied by Westlake and Miller [27]. They reported that the addition of either Nb to Ta or Ta to Nb appreciably increases the terminal solubility of hydrogen (TSH). They concluded that this (TSH) enhancement could be caused by the trapping of hydrogen near substitutional solute sites. They attributed the trapping process to an attractive elastic interaction which had its origin in the atomic size differences in the alloy matrix. Matsumoto and co-workers [28, 29] reported enhancement of terminal solubility in niobium alloys containing Ti and Mo. The P - C - T diagram has been used to determine thermodynamic properties for the niobium-hydrogen system. Various workers examined different temperature ranges, e.g. Albrecht and co-workers [4] from 100 to 900°C, Katz and Gulbran [30] from 225 to 513°C, Veleckis [31] from 350 to 670°C and Kmojathy [13] from 400 to 1000°C. The data for relative partial molal enthalpies, relative partial molal entropies and enthalpies of formation of niobium hydrides are summarized in Table 2 and in Figs 6 - 8 . It is clear that the data obtained by Veleckis [31] and Katz [30] are in a good agreement. It seems that the values obtained by Komjathy [13] are higher; for instance, at H/M = 0 . 3 , Komjathy's [13] data for relative partial molal enthalphies are about 25% higher than Veleckis's [31] and Katz's [30] data. On the other hand, the values

A. Y. ESAYED and D. O. NORTHWOOD

46

Table 2. Data for relativ e partial molal enthalpies (AH), relative partial molal entropies (AS), enthapies of formation (AH) and entropies (AS) of niobium hydrides

H/M

0.05

0.30

0.50

0.60

-AH cal g atom i

-AS eu g atom ~*

-AHf calmo1-1

-ASf cal °C -j mo1-1

8230 8765 8200 10680

7.32 8.16 7.53 10.00

--

--

500

0.37

[33] [34] [5] [14]

10100 10095 9495 13020

13.52 12.38 11.78 16.00

2700 2800 -3650

2.89 1.81 -4.00

[33] [34] [5] [14]

10910 10775 10100 14400

14.39 14.21 13.47 19.00

4790 4880 -5660

6.24 5.66 -6.50

[33] [34] [5] [14]

10900 10960 9900

15.22 15.10 13.76

4790 --

6.24 --

--

--

[33] [34] [5]

325 510

Ref.

0.17 0.30

*eu = c a l mol i °C t

0.55

0.45 0.35 0.35

"~ I I

9000

14000

0.25

0.15

- A ~ ' c a l ( g m a t o m ) -1

Fig. 6. Hydrogen to metal ratio vs - A H for N b - H system. Data taken from O [33], • [5], • [34]. • [14].

0.05

I

1000

obtained by Albrecht [4] are a little smaller. These differences might have arisen from the different methods used to determine the t h e r m o d y n a m i c data. For instance, Veleckis [31] has derived a semi-empirical formula for determining the t h e r m o d y n a m i c quantities.

4.2. Vanadium- hydrogen system Kofstad and Wallace [32] as well Veleckis [31] have used the P - C - T method to obtain the t h e r m o d y n a m i c properties for the V - H system. A comparison o f their results is given in Table 3 and presented in graphical form in Figs 9 - 11. There is a good agreement between the two sets o f t h e r m o d y n a m i c data. Lynch [33] studied vanadium rich alloys o f the V-Ti-Fe-H system and measured the t h e r m o d y n a m i c

I

2000

I

3000

I

4000

I

5000

r

6000

- A H f cal (grn a t o m ) 1

Fig. 7. Hydrogen to metal ratio vs enthalpy of formation for N b - H system. Data taken from O [33], • [34], [] [14].

data o f hydride formation as a function o f alloy composition. This t h e r m o d y n a m i c data is summarized in Table 4. It appears that the ~x ~ /3 transition is suppressed to below room temperature in V - T i alloys when the Ti content is in excess o f 20 at% [ 1 ]. H o w e v e r , Fe has the opposite effect, i.e. it raises the critical temperature o f the tx ,~ /3 transition and this transition again takes place at room temperature when Fe is added to V - T i alloys containing Ti > 20% [I].

GROUP V METAL HYDRIDES

47

Table 3. Thermodynamic properties, relative partial molal enthalpies (AH), relative partial molal entropies (AS), enthalpies of formation (2~H) and entropies of formation (AS0 for the vanadium-hydrogen system H at%

- m/-/rt cal g atom -I

-- ASH eu g atom i

-- mnf cal g atom i

-- msf cal g atom -I

Ref.

5

7790 7316

8.75 8.07

407 374

0.36 0.31

[36] [35]

15

8170 7914

11.21 10.89

1382 1320

1.60 1.51

[36] [35]

25

8600 8446

12.92 12.78

2693 2601

3.49 3.37

[36] [35]

33.3

8690 8400

14.59 14.25

4137 4017

5.77 5.62

[36] [35]

0.55

A

jf

35

0.45

•

25

0.35

/

-~. 0.25

•

E (3) O3 "O

-r 5

0.15

jj 8000 '

7000

90'0 0

-AHHcal (gin atom) -1

0.05 7

•

Fig. 9. Hydrogen composition vs -AHH of V - H system. Data taken from O [361, • [351. I

11

I

15

I

19

-As eu (gmatom)-1 Fig. 8. Hydrogento metalratiovsentropyfor Nb-H system.Data takenfrom O [34], • [33], [] [5], • [14].

oxygen slows the rate of the reaction considerably. This is apparently a result of trapping of hydrogen in O - H complexes. In addition they found deuterides form more slowly than hydrides because of the lower rate of diffusion of deuterium in the metal compared to hydrogen.

4.3. Tantalum-hydrogen system Beleckis [31], Kofstad and Wallace [32], and Mallett and Koehl [34] all used the P - C - T technique to determine the thermodynamic properties of the T a - H system. Their results are shown in Table 5 and in Figs 12 and 13. It can be seen that the partial molal enthalpy results of Veleckis [31] and Mallett and Koehl [34] increase with hydrogen content. The values increase by about 10% in going from 5 at% hydrogen to 25 at% hydrogen. However, the partial molal enthalpies obtained by Kofstad and Wallace [32] are almost constant (there is a small change of about 4-2%) with changing hydrogen content. Lecocq and Wert [35] found that the rate of formation of hydrides in tantalum is markedly dependent on the rate of diffusion of hydrogen in the metal. They also found that

A

N

35f 25

J

J

o

g ~ i

15 5i

1000

2 0

3~

4000

-AHfcal (gmatom)1 Fig 10. Hydrogen composition vs --AHf of V - H system. Data taken from O [36], • [35].

48

A. Y. ESAYED and D. O. NORTHWOOD Table 4. Thermodynamics parameters of the /3 ~ 7 transition in the V - T i - F e - H Desorption X

(kJ m o l t

H2)

system [ 1 ]

Absorption

(J mol t H2 - °C t)

(kJ tool i H2)

(J mol -I H2 - °C 1)

55.20 ± 2.0 48.85 ± 2.10 44.85 ± 1.40

166.1 ~ 5.9 151.4 ± 6.1 150.9 ± 4.1

51.68 ± 2.99

163.4 ± 8.9

(V09Ti0.1)l-~Fe~ 0 0.01 0.02 0.05 0.075

51.79 ± 49.22 ± 47.90 • 43.20 ± 40.00±

0.36 0.60 0.70 2.10 2.70

149.4 146.7 145.1 139.6 136.0

± ± ~ ± ±

1.0 1.7 2.0 6.2 8.1

(V0.sTi02)t-.~Fe~ 0 0.02 0.10

48.1 55.1 49.91 ± 1.50

140.4 151.6 ± 4.3

(V075TiO.25)'l ~Fe~ 0.02

63.5

X = alloy composition (0.00-0.10). AH = enthalpy. AS = entropy.

/// Ta - H s y s t e m

5. H Y S T E R E S I S H y s t e r e s i s is o b s e r v e d in nearly all m e t a l - h y d r o g e n s y s t e m s , d u r i n g h y d r i d e f o r m a t i o n and d e c o m p o s i t i o n . E x p e r i m e n t a l l y it is f o u n d that the e q u i l i b r i u m h y d r o g e n p r e s s u r e for h y d r i d e f o r m a t i o n (Pr) are h i g h e r than the p r e s s u r e on desorption (Pd). H y s t e r e s i s is o f c o n s i d e r a b l e practical i m p o r t a n c e since it reduces the efficiency o f h y d r o g e n (energy) storage s y s t e m s . M a n y e x p l a n a t i o n s h a v e been p r o p o s e d but as yet there is no one theory that has g a i n e d general acceptance. O n e o f the better k n o w n theories is that originally p r o p o s e d by U b b e l o h d e [36]

25 o~

O~

o

"O

\

--r 5

7000

8000 '

90'0 0

10,000 '

-AHf cal (gm atom) t

35

*--

Fig. 12. Hydrogen composition vs - A H for Ta-system. Data taken from O [36], [40]. • [39].

25 Ta - H

g

35

"O

% v 25

I

c13)

5 I

1

I

2-&Sf

r

I

I

3

4

5

i

i

i

5

10

15

-AS H e u (gin atom) q Fig. 11. • O Hydrogen composition vs - A S for V - H system. Data taken from O [36], • [35]. • [] Hydrogen vs - A S f for the same system. Data taken from [36]. [35].

.i

o15 >.

I 5 i

i

9

11

r

13

-AS eu (gm atom) q Fig. 13. Hydrogen composition vs --AS for Ta system. Data taken from O [36], [] [40], • [39].

GROUP V METAL HYDRIDES Table 5. Relative partial molal enthalpies (z~H) and relative partial molal entropies (AS) for the T a - H system Hydrogen at%

- AH (cal g atom -])

-- AS (eu g atom I)

8360 7970 9500

8.20 7.70 9.90

[361

5

10

8660 8280 9400

9.93 9.40 11.10

[36] [40] [39]

15

8830 8840 9500

10.98 11.00 12.10

[36] [40] [30]

25

9320 8690 9700

13.05 12.00 13.60

[36] [40] [39]

Ref.

[401

[39]

over 55 years ago for hysteresis in P d - H systems. He reported that for these systems the free energy of the hydride phase, /3-phase, is not only determined by the temperature, pressure and composition variables but needs other variables such as mechanical strain for its full determination. He argued that in this way it is possible to have more than one pressure for the same concentration of hydrogen. The metallic hydride phase is assumed to be both disordered and strained, thus giving rise to a greater plateau pressure for hydride formation than for hydride decomposition. He then suggested that the phase rule has to be modified for such systems and an additional variable such as strain has to be added. Using a statistical mechanical approach in which the interaction energy between hydrogen atoms in the lattice is taken into account, Lacher [37] derived a theory which assigned metastable states to the ~- and /3-phase. This resulted in oversaturation of the s-phase on absorption and undersaturation of the i3-phase on desorption. Everett and Nordon [38] applied the different parameters of Lacher's [37] statistical mechanical theory to the formation and decomposition plateau branches. These different parameters were suggested to result from the different sites of stress of dilut c~-phase in the hydride phase. Wanger [39] suggested that hysteresis arises from the fact that the metal chemical potentials are not equal in the c~- and 13-phases because the metal atoms are not mobile at the temperatures at which experiments are carried out. Libowitz et al. [40] proposed a theory for hysteresis based on the defects in the lattice. They assumed that there are non-stoichiometric vacancies in the lattice. When the hydrogen is removed from the stoichiometric hydride, the hydrogen vacancies are formed and as a result the hydride becomes non-stoichiometric. At the composition where the lattice becomes saturated with vacancies, further withdrawing of hydrogen causes the lattice to break down, thus forming a two-phase system. Therefore, the plateau pressure is

49

actually the equilibrium pressure of the non-stoichiometric hydride. Flotow [41 ] proposed that metal first formed on desorption of hydrogen from a hydride would be microcrystalline, while the metal absorbing hydrogen would be in a macrocrystalline form due to sintering. The difference in sizes of the crystallites would give different enthalpies of reaction, thus causing different hydrogen pressures on absorption and desorption. Schottus and Hall [42] related hysteresis to plastic deformation during hydride formation. They assumed that strain was absent during the hydride decomposition and as such the decomposition plateau pressure corresponds to true equilibrium. This is probably not a correct assumption, and in fact Flanagan and Clewely [43] have proposed a model to explain hysteresis which is based on the requirement for dislocation production during both hydride formation and hydride decomposition. Kuijpers and van Mal [44] proposed an explanation of hysteresis observed in L a N i s - H and S m C o s - H systems. They argued that the magnitude of the hysteresis was due to larger volume expansion which accompanies the formation of the hydride. Lundin and Lynch [45] proposed an atomic model to explain hysteresis in intermetallic compounds. It focuses on the effects of hydrogen atoms occluded in the interstitial sites in the lattice and the resulting strains imposed. They found a correlation between the standard free energy of formation of the hydride and the size of the interstitial hole in the metal lattice. A critical review of the different theories for hysteresis has been given elsewhere [46]. Birnbaum et al. [13] observed hysteresis in the solves behaviour for the N b - H system. They have concluded that the thermal hysteresis was due to plastic deformation during both processes (i.e. hydriding and dehydriding). Their model for the N b - H system suggests that neither Pf nor P~ corresponds to equilibrium since non-equilibrium defects are generated along each plateau pressure branch.

! "~ 2.0 ,...,

0.5 Q. c

9

o.1

I

0.02

0.8

112

116

21o

H/M

Fig. 14. Pressure-composition isotherms for V - H system. • absorption, O desorption [55].

50

A.Y. ESAYED and D. O. NORTHWOOD

I0.000

10,000

/.

/

I 1000

13_

1000

I

2

00

lOO

112

i:o

2:0

11o

HIM Fig. 15. P - C i s o t h e r m s

for V , N b l

~.-H systems (x = •

1:4 HIM

A 0.1,

• [] 0.2, • O 0.3). Filled symbols correspond to absorption data and open symbols to desorption data. T = 318 K [4]. Lynch et al. [19] have measured hysteresis for the N b - H system and they found it was greater by a factor of 4 than that of V - H system. Figure 14 shows their results for the P - C isotherm for the monohydride-to-dihydride transition at 313 K, in the V - H system. Esayed and Northwood [3] have determined the hysteresis factor for Vx Nb, x alloys (where x = 0.1, 0.2 and 0.3) and they found the hysteresis effect as described by the ratio (Pf/Pd) for x -- 0.3 is about 60% larger than for x = 0.1 at the same temperature. They also found that the hysteresis effect as given by the ratio (Pf/Pd) decreases with increasing temperature and so does the logarithmic function 1/2RT ln(Pf/Pd), which represents the Gibbs free energy loss per mol of atomic hydrogen in completing a hysteresis loop. For instance the hysteresis factor I / 2 R T ln(Pf/Pd) changes from 3340 J mol -~ H at 318 K to 2450 J m o l - ' H at 373 K for x = 0.3, which is about a 30% change. Figure 15 shows the hysteresis effect at 318 K for the V ~ N b , _ x - H systems where x = 0.1, 0.2 and 0.3 Figure 16 shows the effect of temperature on hysteresis for t h e Wo.iNb0.9. Both Figs 15 and 16 were taken from Ref. [41. Maeland et al. [47] found that alloying Nb with Cr or Fe induced Nb to react rapidly with hydrogen at room temperature and alloying Nb with Ta or Zr showed no effect on the reaction of Nb with hydrogen.

6. KINETICS In this section some theories which explain the slow reaction rate of hydrogen with bcc metals are discussed.

Fig. 16. P - C isotherms for V0 iNbo.9-H system. • o 813 K, • [] 353 K, • A 373 K. Filled symbols correspond to absorption data and open symbols to desorption data [4]. Pick and co-workers [48, 49] proposed a model to explain the slow reaction rates of hydrogen with bcc metals. This model is based on the idea that there is a surface barrier which is the result of a strongly bound surface state for hydrogen. This surface state is of lower energy than the bulk state and therefore is filled preferentially over the bulk. The model shows that even though there is a low concentration in the bulk, there can be a high concentration on the surface which reduces the uptake rate of hydrogen. Lagos and co-workers [ 5 0 - 5 2 ] proposed a model based on hydrogen trapped at the subsurface which acts as a barrier (or valve) which controls the passage of further hydrogen into the bulk at low temperature. They reported [52] that at low temperature (300°C) the surface valve is closed and the bulk and the surface are completely decoupled. At higher temperatures the valve opens up slightly and therefore equilibrium is obtained between the surface and the bulk. Around 600°C the surface valve is completely open. This model explains why the bcc metals do not react with hydrogen at room temperature and why it is important to increase the temperature to start the reaction. There is some experimental evidence [53] for subsurface states (as proposed by Lagos et al. [ 5 0 - 5 2 ] ) in niobium. Examination of the surface of a niobium samples after cleaning under ultra-high vacuum prior to exposure to hydrogen showed electron energy loss peaks at room temperature of 0.117 and 0.130 eV. These peaks are assumed to be due to subsurface hydrogen. Fang et al. [ 5 4 ] , who used an angle-resolved photoemission technique, confirmed the existence of these subsurface states. Libowitz and Maeland [ 1 ] explained the rapid reaction of the bcc solid solutions with hydrogen in terms of the sub-

GROUP V METAL HYDRIDES Table 6. Free energies of formation on the group V transition metal-hydrogen systems (AGf) cal mol- i HIM

NB-H

0.1 0.15 0.20 0.25 0.30 0.35

-400 -475 -520 -555 -570 -600

Ta-H -340 -410 -480 -510 -520 -550

V-H -245 -285 -320 -340 -360 -380

*The free energy (AGf) calculated from the entropy and enthalpy data by using the following equation: /XGf = ( A H -

TAS)/Xr,.

surface model. In this explanation the presence of smaller metal (solute) atoms induces strains in the parent metal bcc lattice such that the subsurface trapping energy of the hydrogen is reduced, The surface barrier to the transfer of hydrogen at or below room temperature is thus removed. Rapid reaction with hydrogen occurs, since the diffusion of hydrogen in the bulk material is very fast for bcc metals. Libowitz and Maeland [ 1] also suggested that another possible explanation for the rapid reaction of bcc solid solutions with hydrogen is that the formation of the oxide layer on the surface is modified somewhat by the alloying (solute) metal such that hydrogen can permeate the modified oxide layer more rapidly. These models are summarized in Table 7.

51 7. S U M M A R Y

A review of the crystallography shows that the vanadium and tantalum hydrides have a body centred tetragonal (bct) structure at low hydrogen contents close to the two-phase (t~ +/3) region. In the case of tantalum the hydrides become orthorhombic as the hydrogen contents increases. For niobium the B-phase has an orthorhombic structure at low temperature and has a composition of Nb0.8-0.9. The dihydrides of niobium and vanadium have been shown to be FCC. No dihydride has been reported for tantalum. From the thermodynamic data presented by Mallett and Koehl [ 3 4 ] , Kofstad and Wallace [ 3 2 ] , and Veleckis [31 ], it appears that the n i o b i u m - h y d r o g e n solid solution is more stable than V and Ta hydrogen solid solution. The difference in stabilities is shown to be due to differences in enthalpy values (AH), since the entropies (AS) are almost the same for all systems. Therefore, the hydrogen-to-metal bonding is the strongest in niobium and the weakest in vanadium. There is a decrease in partial entropy with increasing composition (hydrogen content) in each case, which is expected for all solid solutions. The values of partial enthalpy of hydrogen become more negative as the hydrogen content increases, which is an indication that the solution of hydrogen becomes more exothermic. A number of models have been put forward to explain the slow reaction rate of hydrogen with these bcc metals. Pick and Green [49, 50] attributed this phenomenon to the prescence of an intrinsic barrier which is the result of a strongly bound surface state for hydrogen. Lagos et al. [ 5 1 - 5 3 ] atributed the slow reaction rate of hydrogen with bcc metals to the existence of a subsurface " v a l v e " which controls the transfer of hydrogen through the surface. Libowitz and Maeland [1] explained the fast reaction of

Table 7. Summary of different models for reaction rates of hydrogen with bcc metals, Nb, V and Ta Author

Model based on

Ref.

Pick and co-workers

There is an intrinisc barrier to hydrogen absorption. The model did not include the deep subsurface potential. Also did not consider the temperature dependence of the initial sticking coefficient.

[48, 49]

Lagos and co-workers

Existence of a subsurface 'valve' which controls the transfer of hydrogen through the surface. This explains in a natural way the full kinetics of absorption, including the time and temperature dependence of hydrogen absorption. If this model is correct, the energy loss peaks seen in HREELS study of niobium [52] should also appear when alloying Nb with metals which do not cause a fast reaction such as T and Ti or Zr. 1. Smaller metal (solute) induced strain, as a result of the subsurface trapping energy being reduced. 2. Formation of oxide layer on the surface is modified and the hydrogen can inter the modified oxide layer on the metal surface.

[50-52]

Libowitz and co-workers

[1 ]

52

A . Y . ESAYED and D. O. NORTHWOOD

h y d r o g e n with bcc solid solutions as being due to the s m a l l e r metal a t o m s i n d u c i n g strains in the parent metal bcc lattice. Libowitz and M a e l a n d [ 1] also p r e s e n t e d a n o t h e r possible e x p l a n a t i o n for the a b o v e p h e n o m e n a n a m e l y that the o x i d e layer on the s u r f a c e is modified by the solute a t o m s and thus the h y d r o g e n can m o r e easily enter t h r o u g h this m o d i f i e d o x i d e layer. Both alloy c o m p o s i t i o n and t e m p e r a t u r e effect the h y s t e r e s i s in this group. T h e r e is a d e c r e a s e in the d e g r e e o f h y s t e r e s i s as the t e m p e r a t u r e increases. T h e r e are no reports for the effect o f h y d r i d e f o r m a t i o n a n d d e c o m p o s i t i o n cycling on h y s t e r e s i s and plateau p r e s s u r e in this g r o u p . E x p e r i m e n t a l w o r k is underw a y at the U n i v e r s i t y o f W i n d s o r to provide i n f o r m a t i o n on the effect o f h y d r i d e f o r m a t i o n and d e c o m p o s i t i o n p r e s s u r e c y c l i n g on h y s t e r e s i s for the N b - V - h y d r o g e n system. S o m e o f this w o r k will be p u b l i s h e d soon. Acknowledgement. The authors are indebted to the Natural Science and Engineering Research of Canada for financial support of this work through an Operating Grant (A4391) to Prof. D. O. Northwood. REFERENCES 1. G. G. Libowtiz and A. J. Maeland, Mater Sci Forum 177, 31 (1989). 2. E. Formm and H. Uchida, J. Less-Common Metals 66, 77 ( 1978) 3. A. Y. Asayed and D. O. Northwood, Int. J. Hydrogen Energy 16, 687 (1991). 4. W. M. Albrecht, W. D. Goode and M. Mallett, J. Electrochem. Soc. 106, 981 (1959). 5. M. W. Mallett and B. G. Koehl, J. Electrochem. Soc. 109, 968 (1962). 6. W. M. Mueller, J. P. Blackledge and G. G. Libowitz, Metal Hydrides. Academic Press, New York (1968). 7. G. Brauer and H. Mueller, Angew. Chem. 70, 53 (1957). 8. G. Brauer and H. Mueller, J. lnorg, Nucl. Chem. 17, 102 (1961). 9. G. Brauer and R. Hermann, Z. Anorg. All. Chem. 274, 11 (1953). 10. A. J. Maeland, T. R. P. Gibb and D. P. Schumacher, J. Am. Chem. Soc. 83, 3728 (1961). I1. G. V. Samsonov and M. M. Antonova, Zh. Fiz. Khim. 35, 900 (1961). 12. H. K. Birnbaum, M. L. Grossbeck and M. Amano, J. LessCommon Metals 49, 357 (1976). 13. S. Komjathy, J. Less-Common Metals 2, 466 (1960). 14. J. J. Reilly and R. H. Wiswall, lnorg. Chem. 9, 1678 (1970). 15. H. W. Paxton, J. M. Sheehand and W. J. Boleyak, Trans AIME. 215, 725 (1959). 16. C. Wainwright, Bull. Inst. Metals 48, 68 (1958). 17. C. Wainwright, A. J. Cook and B. E. Hooking, J. LessCommon Metals 6, 362 (1964). 18. D. P. Smith, Hydrogen in Metals. University of Chicago Press, Chicago (1948). 19. J. F. Lynch, G. G. Libowitz and A. J. Maeland, J. LessCommon Metals 103, 117 (1984).

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