Studies on the Synergetic Effects in Multi-Phase Metal Hydride Alloys

Abstract

The electrochemical reactions of multi-phase metal hydride (MH) alloys were studied using a series of Laves phase-related body-centered-cubic (BCC) Ti_15.6Zr_2.1V₄₃Cr_11.2Mn_6.9Co_1.4Ni_18.5Al_0.3X (X = V, B, Mg, Y, Zr, Nb, Mo, La, and Nd) alloys. These alloys are composed of BCC (major), TiNi (major), C14 (minor), and Ti₂Ni (minor) phases. The BCC phase was found to be responsible for the visible equilibrium pressure plateau between 0.1 MPa and 1 MPa. The plateaus belonging to the other phases occurred below 0.005 MPa. Due to the synergetic effects of other non-BCC phases, the body-centered-tetragonal (BCT) intermediate step is skipped and the face-centered-cubic (FCC) hydride phase is formed directly. During hydrogenation in both gaseous phase and electrochemistry, the non-BCC phases were first charged to completion, followed by charging of the BCC phase. In the multi-phase system, the side with a higher work function along the grain boundary is believed to be the first region that becomes hydrogenated and will not be fully dehydrided after 8 h in vacuum at 300 °C. While there is a large step at approximately 50% of the maximum hydrogen storage for the equilibrium pressure measured in gaseous phase, the charge/discharge curves measured electrochemically are very smooth, indicating a synergetic effect between BCC and non-BCC phases in the presence of voltage and charge non-neutrality. Compared to the non-BCC phases, the C14 phase benefits while the TiNi phase deteriorates the high-rate dischargeability (HRD) of the alloys. These synergetic effects are explained by the preoccupied hydrogen sites on the side of the hydrogen storage phase near the grain boundary.

1. Introduction

The synergetic effects in multi-phase metal hydride (MH) alloys refer to the presence of microsegregated secondary phases occurring in the melted alloys that effectively provide beneficial effects [1]. Some examples of synergetic effects in the gaseous phase and electrochemical environment are summarized in

Table 1. Examples of synergetic effects in multi-phase metal hydride (MH) alloys. GP and EC denote gaseous phase and electrochemistry experiments, respectively. HRD: high-rate dischargeability.

Main Phase	Secondary Phase	Main Improvement	Environment	Reference
C14	Zr7Ni10	Capacity and activation *	EC	[10]
C14/C15	Zr7Ni10, ZrNi	Capacity and activation	EC	[11]
C15	Zr7Ni10	HRD	EC	[12]
C15	Zr7Ni10	Capacity	EC	[13]
MgNi	Ti	Cycle stability	EC	[14]
C14/C15	Zr7Ni10 and TiNi	HRD	EC	[15]
C14/C15	Zr7Ni10 and Zr9Ni11	Capacity and HRD	EC	[1]
Zr8Ni21	Zr7Ni10, Zr9Ni11	Capacity and HRD	EC	[16]
C14	Zr8Ni21	Activation, bulk diffusion, cycle stability	EC	[17,18]
Zr7Ni10	C15	HRD	EC	[19]
C14/C15	Zr7Ni10 and TiNi	Capacity and reversibility	GP	[20]
C14/C15	Zr9Ni11	Activation, HRD, charge retention, and cycle stability	EC	[21]
C14	Zr9Ni11 and TiNi	Capacity	GP	[21]
C14	Zr9Ni11 and TiNi	Capacity, charge retention, and cycle stability	EC	[21]
C14	Zr7Ni10 and ZrNi	HRD	EC	[21]
C14/C15	Zr7Ni10 and TiNi	HRD	EC	[22]
Mg	Mg2Ni	Desorption kinetics	GP	[23]
AB5	AlMnNi2	Capacity and HRD	EC	[24]
BCC, C14	ZrNi	Capacity, activation, and cycle stability	EC	[25]
Zr2Ni7	Zr7Ni0	Capacity	GP, EC	[26]
Zr7Ni0	Zr8Ni21	Capacity	EC	[26]
Zr2Ni7	ZrNi3, ZrNi5, VNi2, VNi3	Capacity	EC	[9]
Zr2Ni7	ZrNi3 and ZrNi5	Capacity	EC	[8]
NdNi5	Nd2Ni7	HRD	EC	[27]
LiBH4	Fluorographite	Desorption kinetics	GP	[28]
Mg(BH4)2	Fluorographite	Desorption kinetics	GP	[29]
MgH2	In, TiMn2 additives	Desorption kinetics	GP	[30]

* indicates that the authors attributed improvement to micro-cracking at the surface instead of to synergetic effects.

. In general, the synergetic effects in the gaseous phase hydrogen storage can improve the storage capacity and reversibility, and are characterized by a continuous transition in the pressure-concentration-temperature (PCT) isotherm, from the plateau pressure corresponding to the phase with a stronger metal–hydrogen (M–H) bond strength to the phase with a weaker M–H bond strength (Figure 7b in [2]). The interface region between two phases is considered to be critical for the synergy to take place. Transmission electron microscopy (TEM) studies have demonstrated the interface between the main C14 phase and other secondary phases (C15, Zr₇Ni₁₀, etc.) are clean [3,4], and strong crystallographic orientation alignment can be established by electron beam back-scattering diffraction pattern studies [5,6]. Some synergetic effects in the electrochemical environment are similar to those in the gaseous phase, i.e., improvement in capacity and hydrogen absorption/desorption rate. Interestingly, the properties established by the scope of the gaseous phase can be further enhanced by synergetic effects in the electrochemical environment. For example, it has been found that the electrochemical discharge capacities of certain multi-phase MH alloys can be increased substantially by lowering the equivalent plateau pressure (observed through open-circuit voltage (OCV) changes during charge/discharge [7]) through synergetic effects [8,9]. Therefore, a systematic study in the difference between the synergetic effects in the two environments is of significant interest.

Laves phase-related body-centered-cubic (BCC) solid solution alloys were chosen for this study. By combining the high-capacity main storage BCC phase with catalytic phases, such as C14, TiNi, and Ti₂Ni, one alloy in the family demonstrated a 30% increase in capacity when discharged at a current density of 100 mA·g^‒1, which is adequate for electric vehicle applications [31]. The development of this family of alloys proceeded in stages and the X-ray diffraction (XRD) and PCT analyses demonstrated strong evidence of synergetic effects between the main storage phase and the catalytic phase [32,33,34,35]. With a base alloy of P17 (Ti_15.6Zr_2.1V₄₄Cr_11.2Mn_6.9Co_1.4Ni_18.5Al_0.3), the V/Ni content was adjusted to increase the high-rate dischargeability (HRD) [36]. The synergetic effects in an electrochemical environment were studied in detail using this same base alloy, P17, with substitutions of other A-site atoms (B, Mg, Y, Zr, Nb, Mo, La, and Nd) and presented here.

2. Experimental Setup

The alloy samples were prepared using an arc-melting technique. The melting was performed in an Ar environment with an average alloy weight of 12 g. The chemical composition of the ingot, compared to the ratios in the raw materials, was determined with a Varian Liberty 100 inductively coupled plasma-optical emission spectrometer (ICP-OES, Agilent Technologies, Santa Clara, CA, USA). The microstructures of the as-prepared samples were examined by a Philips X’Pert Pro X-ray diffractometer (XRD, Philips, Amsterdam, The Netherlands) and a JEOL-JSM6320F scanning electron microscope (SEM, JEOL, Tokyo, Japan) equipped with energy-dispersive spectroscopy (EDS). The PCT analysis was performed with a Suzuki-Shokan multi-channel PCT (Suzuki Shokan, Tokyo, Japan) system. The MH alloys were compacted on Ni mesh to achieve the negative electrode, and the battery cells were made by using a standard Ni(OH)₂-positive electrode and 30 wt% KOH electrolyte. The electrochemical testing of the battery cells was performed at room temperature (RT) using a CTE MCL2 Mini (Chen Tech Electric MFG. Co., Ltd., New Taipei, Taiwan) cell test system.

3. Results

3.1. X-Ray Diffraction Structure Analysis

Eight alloys with the targeted compositions listed in

Table 2. Design compositions (bold) and inductively coupled plasma (ICP) results for alloys in this study in at %.

Alloy	Source	Ti	Zr	V	Cr	Mn	Co	Ni	Al	X	B/A
P17	Design	15.6	2.1	44	11.2	6.9	1.4	18.5	0.3	0	4.65
	ICP	15.6	2	44.1	11.3	6.4	1.4	18.9	0.3	0	4.68
P38 (X = B)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15.7	2.2	42.7	10.6	6.7	1.4	19.3	0.3	1.1	4.26
P39 (X = Mg)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15.5	2.2	44.5	10.6	7.1	1.4	18.3	0.5	0	4.65
P40 (X = Y)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15.1	2	44.9	11.4	5.4	1.4	18.7	0.3	0.8	4.59
P41 (X = Nb)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15.8	2.1	41.2	11.7	7.4	1.5	19.2	0.4	0.8	4.35
P42 (X = Mo)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15	2	43.1	11	7.7	1.4	18.5	0.4	1	4.56
P43 (X = La)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	16.1	2.2	40.9	11.8	7.4	1.5	19.3	0.4	0.4	4.34
P44 (X = Nd)	Design	15.6	2.1	43	11.2	6.9	1.4	18.5	0.3	1	4.35
	ICP	15.9	2.2	43	9.8	6.8	1.4	19.6	0.3	1	4.24

were prepared by arc melting. The base alloy, P17, originated from a composition optimization study [35]. Other alloys were derivatives of this base alloy with 1 at% replacement aimed at the A-site element (B, Mg, Y, Nb, Mo, La, and Nd). The chemical composition of the as-prepared ingot was verified by ICP and shows a large discrepancy only in alloy P39, where the added Mg was not found in the ingot due to loss from evaporation during melting. Mg has a high vapor pressure at temperatures near the melting point of the alloy and low solubility in the AB₂ phase [37,38]. The content of La in the final ingot of P43 was less than 50% of the targeted content due to the high chemical reactivity of La metal with the residual oxygen and the formation of oxide slag during melting. The B/A ratios for the ingots ranged from 4.24 to 4.68, assuming that Ti, Zr, and the additive elements are A-site atoms and the remaining constituent elements are B-site atoms. The XRD patterns of the alloys are shown in Figure 1. The three major peaks detected in all alloys belong to a BCC structure. All the minor peaks can be attributed to a C14 Laves phase. A TiNi phase with a B2 structure (a close derivative of a BCC structure) and a slightly larger lattice constant were separated from the BCC phase using software deconvolution (JADE 9, Christchurch, New Zealand).

The lattice constants, crystallite size, and phase abundances calculated from the XRD pattern are listed in

Table 3. Lattice parameters, unit cell volumes, phase abundances, and crystallite sizes of phases derived from XRD analysis. XS denotes crystallite size.

Alloys	BCC			C14						TiNi
	a (Å)	XS (Å)	Abundance (wt%)	a (Å)	c (Å)	c/a	Unit Cell Volume (Å3)	XS (Å)	Abundance (wt%)	a (Å)	XS (Å)	Abundance (wt%)
P17	2.977	171	52.8	4.912	8.010	1.631	167.3	215	13.2	2.993	170	34.0
P38	2.961	181	60.5	4.875	7.998	1.641	164.6	631	2.6	2.974	65	36.9
P39	2.968	260	53.0	4.892	8.019	1.639	166.2	322	7.0	2.982	89	39.9
P40	2.974	343	31.8	4.929	8.030	1.629	169.0	414	6.1	2.986	102	62.1
P41	2.971	187	50.4	4.906	7.991	1.629	166.6	309	20.7	2.986	41	28.9
P42	2.973	221	58.4	4.893	8.013	1.638	166.1	808	0.8	2.987	74	40.8
P43	2.962	262	38.9	4.899	8.003	1.634	166.3	240	9.5	2.976	85	51.6
P44	2.973	228	41.0	4.924	8.007	1.626	168.1	278	6.3	2.988	97	52.7

.

The lattice constants a from the three constituent phases are plotted against the metallic radii of the additive elements in the Laves-phase alloy [39] in Figure 2a. At a glance, the lattice constants from all three phases follow the same trend: a rapid increase followed by a decrease with increasing radii of the additive. However, the later EDS analysis indicates that most of the additives have zero or limited solubility in these three main phases. Therefore, the changes in lattice constant are tied to other characteristics of this group of alloys. For example, the lattice constants show a reasonably consistent trend of decreasing with increasing Mn-content (with a relatively small radius) in the BCC phase, as shown in Figure 2b. The spread of c/a ratio (a parameter that can be used to estimate the preferential occupation site for foreign atoms [36,40]) in the A-atom substitution (1.626–1.641) is approximately the same as the spread in the B-atom substitution study (1.627–1.644) [35]. No preference for occupation site substitution can be deduced in the current study. From the abundances shown in Table 3, BCC (32‒61 wt%) and TiNi (29‒62 wt%) are the two major phases, with C14 as the secondary phase. The crystallite of the secondary phase (C14) is larger than those in the main phases (BCC and TiNi), which has previously been observed in BCC-Laves-related alloys [33,36], but which is not common in multi-phase AB₂ MH alloys [41]. It is interesting to observe that there is a correlation between the abundance and the crystallite size in each phase. For example, as the TiNi phase becomes more popular (larger abundance), its crystallite becomes larger. In the case of BCC and C14 phases, the trend is the opposite (Figure 3).

3.2. Scanning Electron Microscope Phase Analysis

The composition of the constituent phase was studied through a combination of SEM and EDS on a few selected areas. A representative SEM-backscattering electron image (BEI) from each alloy is shown in Figure 4, and the corresponding EDS data are listed in

Table 4. Energy-dispersive spectroscopy (EDS) composition data from select spots in Figure 4. All numbers are in at%. B/A ratios for the C14 phase were calculated assuming V is in the B-site, while those in the TiNi and VNi phases were calculated assuming V in the A-site. Compositions in the main BCC phase are highlighted in bold.

Location	Zr	Ti	V	Cr	Mn	Co	Ni	Al	X	B/A	e/a	Phase
P17-1	16.2	16.3	7.8	4.4	5.9	1.8	47.5	0.1	0.0	2.08	7.28	C14
P17-2	4.7	36.4	7.8	0.8	3.6	2.0	44.4	0.3	0.0	1.04	-	TiNi
P17-3	6.2	48.4	11.7	1.2	2.3	1.8	28.3	0.1	0.0	0.51	-	Ti2Ni
P17-4	0.1	5.5	63.5	16.9	7.4	0.8	5.6	0.2	0.0	-	-	BCC
P38-1	9.9	22.0	20.3	4.7	6.6	2.0	33.9	0.6	0.0	2.13	6.62	C14
P38-2	4.4	37.0	6.4	1.1	2.8	2.7	44.8	0.8	0.0	1.09	-	TiNi
P38-3	5.5	45.4	15.4	2.4	2.8	1.8	26.4	0.3	0.0	0.51	-	Ti2Ni
P38-4	0.1	5.2	62.8	17.5	8.4	0.8	5.0	0.2	0.0	-	-	BCC
P38-5	0.1	13.1	73.4	8.8	3.2	0.2	1.1	0.1	0.0	6.58	-	Oxide
P39-1	9.7	22.8	20.0	4.8	7.1	2.2	32.5	1.0	0.0	2.08	6.56	C14
P39-2	3.8	33.9	11.9	2.4	3.4	2.8	40.6	1.1	0.0	1.01	-	TiNi
P39-3	6.3	48.5	11.4	1.4	2.8	1.8	27.5	0.3	0.0	0.51	-	Ti2Ni
P39-4	0.1	5.7	61.7	18.5	7.8	0.8	5.1	0.3	0.0	-	-	BCC
P40-1	11.8	21.3	20.4	3.3	4.8	1.6	36.2	0.7	0.0	2.02	6.66	C14
P40-2	5.2	37.9	5.5	0.9	2.0	2.8	45.2	0.4	0.0	1.06	-	TiNi
P40-3	0.1	4.6	66.2	17.1	7.2	0.8	3.7	0.2	0.0	-	-	BCC
P40-4	2.0	0.8	1.4	0.5	0.3	0.0	53.4	0.8	40.8	1.22	-	YNi
P40-5	6.6	84.9	5.7	0.5	0.4	0.1	1.7	0.2	0.0	-	-	TiO2
P40-6	0.0	1.4	1.7	0.4	0.0	0.0	2.3	1.3	92.9	-	-	Y
P41-1	6.1	22.9	21.1	5.9	7.4	2.0	31.2	0.5	2.8	2.14	6.54	C14
P41-2	2.8	37.5	7.2	1.4	3.3	2.7	43.8	0.8	0.5	1.08	-	TiNi
P41-3	4.2	47.9	13.3	2.2	2.7	1.8	27.0	0.3	0.6	0.52	-	Ti2Ni
P41-4	0.1	5.2	62.9	17.7	7.1	0.7	5.7	0.2	0.5	-	-	BCC
P42-1	10.3	21.9	20.0	4.4	6.4	2.0	34.3	0.6	0.1	2.10	6.63	C14
P42-2	4.5	37.0	6.3	0.9	2.6	2.7	45.2	0.8	0.0	1.09	-	TiNi
P42-3	6.2	45.5	14.8	2.3	2.5	1.6	26.8	0.3	0.0	0.51	-	Ti2Ni
P42-4	0.4	7.0	58.6	15.4	7.9	1.1	8.0	0.2	1.4	-	-	BCC
P43-1	9.2	23.5	20.6	5.3	7.3	2.1	31.3	0.7	0.0	2.06	6.51	C14
P43-2	4.1	38.5	5.9	1.1	2.5	3.0	44.3	0.7	0.0	1.06	-	TiNi
P43-3	0.1	5.9	61.1	18.6	8.1	0.8	5.2	0.2	0.0	-	-	BCC-1
P43-4	1.1	30.2	45.5	12.7	6.3	0.5	3.3	0.1	0.3	-	-	BCC-2
P43-5	0.2	1.3	2.6	0.8	0.0	0.3	38.3	0.1	56.4	0.65	-	LaNi
P44-1	9.9	23.1	20.8	4.4	6.4	2.1	32.8	0.6	0.0	2.03	6.56	C14
P44-2	4.3	38.4	6.0	0.9	2.5	2.8	44.4	0.7	0.0	1.05	-	TiNi
P44-3	0.1	5.8	64.0	16.2		0.8	5.1	0.2	0.0	-	-	BCC
P44-4	0.5	2.9	2.9	0.3	0.9	0.1	46.5	0.2	45.7	0.92	-	NdNi
P44-5	0.1	1.8	4.5	0.7	1.2	0	2	0	89.8	-	-	Nd

. In contrast to the composition information obtained from ICP, EDS results are far less accurate and representative but are still useful for qualitatively studying the composition in each constituent phase. The two main phases, BCC and TiNi, form separate 3D frameworks that interlace with each other (for a schematic drawing, see Figure 4 in [32]). This microstructure suggests an early formation of a V-rich BCC framework, which pushes the Ni and Ti into grain boundaries and later forms the C14 and TiNi phases. According to an earlier TEM study of the C14-predomintaed MH alloy, the C14 phase is solidified before the formation of the TiNi phase [3,4]. The B/A ratio in each spot was calculated assuming B occupies the A-site in the C14 phase [42] but occupies the A-site in the TiNi and Ti₂Ni phases [34]. All the calculated B/A ratios are slightly above their stoichiometric values (hyper-stoichiometry), i.e., 2.0 (C14), 1.0 (TiNi), and 0.5 (Ti₂Ni). The anti-site defect (Ni in the A-site) is energetically more favorable, compared to the vacancy defect [43], which explains the hyper-stoichiometry. The average electron density (e/a), calculated from the number of conduction electrons in the constituent metals [44] in the C14 phase, is lower than the C14/C15 threshold [45] and thus a C14-structure is expected instead of a C15-structure, except for alloy P17. There may be some of the C15 phase in alloy P17, according to its relatively higher e/a ratio. From the EDS study, we found that only additive Nb participated in all three main phases and only Mo was found in the BCC phase. Boron (B) is too light to be detected by EDS, but definitely exists in the alloy, and Mg is missing completely, which can both be confirmed by ICP. Rare earth elements, such as Y, La, and Nd, promote formation of the AB phase and have undetectable solubilities in the BCC, TiNi, and C14 phases. The Mn-content in the C14 phase has been correlated with the lattice constants of the three phases in Figure 2b. Other correlations are less obvious.

3.3. Gaseous Phase Characteristics

The characteristics of gaseous phase hydrogen storage for these eight alloys were studied by PCT analysis and the resulting isotherms, obtained at 30 °C and 60 °C, are shown in Figure 5. Similar to other alloys in the same family [32,33,35,36], the PCT isotherm shows only one plateau in the pressure range of our apparatus (0.001–5 MPa). There is at least one more known plateau below 0.001 MPa. The PCT hysteresis from this family of alloys is much larger than those from the AB₂ [46], AB₅ [47], and A₂B₇ [27] MH alloy families. The gaseous phase properties obtained from the PCT analysis are summarized in

Table 5. Summary of gaseous phase and thermodynamic properties.

Alloy	Desorption Pressure @30 °C (MPa)	Desorption Pressure @60 °C (MPa)	Maximum Capacity @30 °C (wt%)	Reversible Capacity @30 °C (wt%)	PCT Hysteresis @30 °C	PCT Hysteresis @60 °C
P17	0.17	0.53	1.68	1.01	1.06	1.10
P38	0.21	0.65	1.64	1.13	1.14	1.06
P39	0.21	0.71	1.72	1.14	1.16	1.06
P40	0.19	0.60	1.69	0.85	1.22	1.06
P41	0.23	0.76	1.57	1.15	1.13	1.05
P42	0.40	1.29	1.65	1.02	1.13	0.91
P43	0.28	1.05	1.85	1.09	1.23	0.76
P44	0.15	0.47	1.69	0.88	1.27	1.22

. The desorption plateau pressures of the substituted alloys are higher than that in the base alloy (P17), except for Nd-doped P44, which indicates that most of the additives weaken the M–H bond in the hydride. By comparing the capacities of these alloys, we found that while La is beneficial to both maximum and reversible capacities, both B and Nb increase the reversible capacity. The irreversible hydrogen storage capacity (the difference between the maximum and reversible capacities) is proportional to the TiNi phase abundance (shown in Figure 6). The PCT hysteresis originates from the elastic lattice deformation energy needed at the metal (α)‒hydride (β) interface during hydrogen absorption [48,49]. The higher PCT hysteresis in this family of alloys suggests an environment which results in difficulty with respect to expansion inside the alloy. The substituted alloys show larger PCT hysteresis at 30 °C, but smaller PCT hysteresis at 60 °C when compared to those from the base alloy P17.

Thermodynamic properties, including changes in enthalpy (ΔH) and entropy (ΔS), were calculated from the equation for free energy (G):

$$\mathrm{\Delta }G=\mathrm{\Delta }H-T\mathrm{\Delta }S=\mathrm{R}T\mathrm{In}P$$

(1)

where R is the ideal gas constant and T is the absolute temperature. Based on plateau pressures in the absorption, desorption, and halfway point, three sets of ΔH and ΔS were calculated and listed in

Table 6. Thermodynamic properties ∆H and ∆S calculated from the equilibrium pressure plateaus (A: absorption, D: desorption, (A + D)/2:half-point) measured at 30, 45, and 60 °C with error bars of approximately ±3%. −∆H_Calc is the calculated value from the BCC phase composition obtained with SEM-EDS (Table 4) using Equation (3).

Alloy	−∆HA (kJ·mol−1)	−∆SA (J·mol−1·K−1)	−∆HD (kJ·mol−1)	−∆SD (J·mol−1·K−1)	−∆H(A+D)/2 (kJ·mol−1)	−∆S(A+D)/2 (J·mol−1·K−1)	−∆HCalc (kJ·mol−1)
P17	29.2	111	30.9	107	30.4	110	28.6
P38	28.7	111	31.7	111	30.2	111	28.0
P39	29.4	113	33.9	118	31.7	115	28.2
P40	30.0	114	32.2	111	31.1	113	28.4
P41	28.3	111	34.0	119	30.8	114	28.1
P42	29.9	119	32.8	120	31.4	120	29.3
P43	28.0	111	37.2	131	32.4	120	28.3
P44	31.6	118	31.6	108	31.8	113	29.1

. Conventionally, the values obtained using the desorption isotherm (ΔH_D and ΔS_D) are adopted because the hysteresis is mainly due to the energy barrier from lattice distortion at the α‒β interface during hydrogen absorption. In the current study, the ΔS_A value is more consistent, suggesting that the absorption isotherm may be increasingly unaffected, and at least should be included in the report. As can be seen in Table 6, ΔH_D is always lower than ΔH_A, and ΔH_(A+D)/2 falls between the two values. ΔS_D can be higher or lower than ΔS_A, and ΔS_(A+D)/2 also falls in the middle. In our previous publication, we adopted an estimation method for ΔH from the atomic percentage and obtained a ΔH of the three hydride-former atoms, Zr (∆H_h = −163 kJ·mol⁻¹ H₂ for ZrH₂), Ti (∆H_h = −124 kJ·mol⁻¹ H₂ for TiH₂), and V (∆H_h = −34 kJ·mol⁻¹ H₂ for VH₂). This estimation is based on the following assumptions. The complete heat of hydride formation of AB_n alloy should be:

ΔH (AB_nH_2m) = Δ H (AH_m) + ΔH B_nH_m) − ΔH (AB_n) [50]

(2)

However, the ΔH for the alloy AB_n and the non-hydride former atoms are smaller than those in the hydride former atoms (for example, ΔH for VCr₂ and NiH_0.5 are −2 kJ·mol⁻¹ H₂ [51] and −6 kJ·mol⁻¹ H₂ [52], respectively). Therefore, the heat of hydride formation can be estimated by:

ΔH (A1_hA2_kA3_lB_nH₂) = hΔH (A1H₂) + kΔH (A2H₂) + lΔH (A3H₂)

(3)

where h, k, l, and m are the atomic percentages of the hydride former elements A1, A2, and A3 and the non-hyride former element B. According to this calculation, the ΔH for the hydride of the constituent phases of alloy P17 are −28.5 (BCC), −55.4 (TiNi) and −49.3 kJ·mol⁻¹ H₂ (C14). Therefore, the pressure plateaus in the PCT shown in Figure 5 should correspond to the BCC phase, and those from TiNi and C14 phases are below the minimum pressure used in our PCT apparatus (0.001 MPa). This explains why the irrevesible part of the capacity (below 0.001 MPa) has a linear correlation with the abundance of the TiNi phase, as shown in Figure 6.

The ΔHs calculated with Equation (3) and the composition of the BCC phase highlighted in Table 4 for each alloy are listed in the last column of Table 6 and compared to the measured ΔH values in Figure 7a. Except for alloy P42 (Mo), the measured ΔHs, no matter the variation from the absorption (A), desorption isotherms (D) or half-point (half), are lower (more negative) and do not vary much compared to the calculated values. The decrease in the measured ΔH is due to the synergetic effects associated with the other two phases that have a much lower ΔH and will be discussed in the next section of this paper. In addition, the plateau pressure for the PCT and BCC phase lattice constants have been correlated to the calculated ΔH in the BCC phase, and the results are shown in Figure 7b. Except for alloy P42 (Mo), it is logical to connect the larger BCC unit cell volume to both lower plateau pressure and lower (more negative) ΔH, indicating a stronger M–H bond. There is an appreciable amount of Mo in the BCC phase of alloy P42 (1.4 at %). Mo has a relatively large atomic size and contributes to the increase in lattice parameter. However, instead of lowering the ΔH, Mo has a positive contribution (ΔH for MoH_0.5 = +10 kJ·mol⁻¹ H₂ [52]). This explains the abnormal behavior of alloy P42 (Mo) in Figure 7.

3.4. Electrochemical Properties

The electrochemical hydrogen storage properties were determined in a flooded half-cell configuration with a commercially available co-precipitated Ni₉₁Co_4.5Zn_4.5 hydroxide (BASF, Rochester Hills, MI, USA) counter electrode and 30 wt % KOH electrolyte. The pressed electrode was charged with a current density of 100 mA·g⁻¹ for 5 h and then discharged with the same current density to a cut-off voltage of 0.9 V with two more pulls at 24 mA·g⁻¹ and 8 mA·g⁻¹ at the same cut-off voltage. The obtained total capacities and HRDs (as defined by the ratio between capacities obtained from 100 mA·g⁻¹ and the total discharge capacity) from the first 13 cycles are plotted in Figure 8a,b, respectively, to demonstrate the activation behavior of these alloys. Most of the alloys reach their maximum capacity at Cycle 3 or 4, except for alloys P38 (B) and P40 (Y) which require more cycles to stabilize the capacity. It is easier to achieve HRD activation for alloys with higher HRD than those with lower HRD (alloys P40, P38, and P43). The capacities from 100 mA·g⁻¹ and 8 mA·g⁻¹ discharge currents and their ratios for Cycle 5 are listed in

Table 7. Summary of electrochemical properties. Cap₁₀₀ and Cap₈ are discharge capacities measured for discharge currents at 100 mA·g^‒1 and 8 mA·g^‒1, respectively. HRD₁₀₀ is the ratio of capacities measured at 100 mA·g^‒1 and 8 mA·g⁻¹. RT: room temperature; and OCV: open circuit voltage.

Alloy	Cap100 at 5th Cycle (mAh·g−1)	Cap8 at 5th Cycle (mAh·g−1)	HRD100 (Cap100/Cap8)	Activation Cycle to Reach Maximum Capacity	Diffusion Coefficient, D @RT (10−10 cm2·s−1)	Exchange Current Io @RT (mA·g−1)	OCV (V)
P17	374.0	397.4	0.948	7	1.69	36.5	1.302
P38	332.6	399.0	0.834	10	2.24	22.3	1.303
P39	360.3	402.1	0.896	4	1.81	35.6	1.311
P40	300.2	389.9	0.770	7	1.75	15.0	1.258
P41	362.5	392.9	0.922	5	2.49	37.6	1.319
P42	351.5	392.4	0.896	7	2.27	19.6	1.258
P43	332.6	382.1	0.870	5	2.57	19.2	1.237
P44	346.1	399.2	0.868	4	1.99	15.4	1.266

. It is obvious that while the increase in the full capacity is only marginal with B (P38) and Nd (P44), the decrease in high-rate capacity and HRD is significant. The additives chosen for this study do not significantly improve the electrochemical properties of the base alloy (P17). The gaseous phase capacity is converted to the electrochemical capacity following 1 wt% H₂ = 268 mAh g⁻¹ and plotted in Figure 9. As in other MH alloys [53], both the low-rate and high-rate electrochemical capacities fall in between the boundaries set by the maximum and reversible gaseous phase capacities. There is no clear correlation between the gaseous phase maximum storage capacity and the electrochemical low-rate capacity, but there is a strong similarity between the gaseous phase reversible capacity and the electrochemical high-rate capacity, except for the base alloy P17 (V).

In order to further study HRD performance, the bulk diffusion constant (D) and surface exchange current (I₀) of each sample were measured at RT according to previously described procedures [54] and are listed in Table 7. We found that while D values in the modified alloys are higher than the base alloy, the opposite is true for I₀ values (except for alloy P41 (Nb)). The increase in the degree of disorder, achieved by introducing an additional element, facilitates the diffusion of hydrogen in the bulk; the decrease in the V-content with a high leaching rate in KOH [55] impedes the surface reaction, except for substitutions with Nb, which is in the same column of V and is expected to have a high corrosion rate in KOH solution. It is difficult to attribute the large decrease observed in I₀ to a 1% decrease in V-content, which will be discussed in the next session. The deterioration of HRD with various additives is mainly due to the decrease in surface electrochemical reaction activity.

The OCV was measured at a state-of-charge of 50% and results are listed in Table 7. Theoretically, OCV can be related to the equilibrium pressure in the gaseous phase through the Nernst equation:

OCV (vs. NiOOH at 0.36 V) = 1.294 + 0.029logp(H₂) [7]

(4)

However, the plot of the desorption pressure versus OCV, as shown in Figure 10, does not follow the trend where higher pressure corresponds to higher OCV, as indicated in Equation (4). However, the OCV in this study seems to correlate closely with the maximum hydrogen storage in the gaseous phase (Figure 10). The alloys with higher OCV show smaller gaseous phase H-storage capacities. Similar phenomena have been observed previously, but were tied to the change in plateau pressure and strength of M–H bonding [56].

4. Discussion

It is generally believed that in a multi-phase MH system the phase with a lower plateau pressure and more negative ∆H is a main storage phase, which has a larger hydrogen storage capacity (Zr₇Ni₁₀, for example [57]), and the secondary phase with a higher plateau pressure and less negative ∆H is considered to be the catalytic phase (Zr₂Ni₇, for example [57]). In the case of funneling phenomenon, the gaseous phase hydrogen storage is accomplished with the catalytic phase as the necessary funnel to move hydrogen in and out of the alloy [2]. In the case of Laves phase-related BCC MH alloy, the BCC phase is considered to be the main storage phase; although it has a large hydrogen storage capacity, it has limited absorption/desorption kinetics that require an additional catalytic phase to facilitate the hydrogen storage process [58]. Therefore, both C14 and TiNi were considered to be catalytic phases in this family of alloys [32,33,34,35,36]. However, the ∆H values estimated from Equation (3) for the C14 and TiNi phases are much lower than those of the BCC phase, and the visible plateau in the pressure between 0.1 MPa and 1 MPa was assigned to the BCC phase. It is necessary to verify the correctness of such an assignment.

In order to investigate the phase abundance evolution during both gaseous phase and electrochemical hydrogen absorption/desorption, an alloy with a target composition of P17 (Table 2) was reproduced through conventional induction melting. The XRD patterns from the as-prepared samples show an additional Ti₂Ni phase (Figure 11a) which was not observed in the sample prepared by arc melting, due to different cooling rates (Figure 1). The Rietveld refinement indicates percentages of 53.6% BCC, 29.4% TiNi, 10.8% Ti₂Ni, and 6.2% C14 (

Table 8. Summary of the phase abundance of a P17 alloy prepared by induction melting through various gaseous phases and electrochemical hydrogen absorption/desorption processes.

Stage	BCC		C14		TiNi		Ti2Ni
	α	β	α	β	α	β	α	β
Pristine	53.6	-	6.2		29.4	-	10.8
300 °C vacuum	54.0	-	3.0	2.0	30.6	-	4.2	6.2
0.005 MPa	51.7	-	-	4.1	11.3	22.2	3.6	7.1
4 MPa	6.1	46.4	-	3.6	-	34.9	-	9.0
50% SOC	53.4	-	-	5.1	0.7	32.0	-	8.8
100% SOC	29.5	21.1	-	4.7	-	36.5	-	8.2

). The sample was activated in a 4 MPa hydrogen environment first and then degassed under vacuum for 8 h at 300 °C. XRD analysis at this stage (Figure 11b) shows a shift for TiNi peaks to lower angles, which indicates the presence of some hydrogen remaining in the α-TiNi. The Rietveld refinement shows the β-phase (MH) of C14 and Ti₂Ni. There are still MH remaining in the alloy after 8 h of degassing. In order to understand the cause of the remaining hydride after 8 h degassing, a discussion of two dissimilar metals is necessary. When two metallic phases, M1 and M2, with different work functions (differences between the electron potential in vacuum (E_VAC) and the Fermi level (E_F)) are brought into contact (Figure 12b), there will be a small charge transfer that builds a potential (contact potential), preventing electron flow from the metal with a smaller work function (M2) into the metal with a larger work function (M1). During the initial hydrogenation, the neutral hydrogen (proton plus a nearby electron) will reside on the M1 side of the boundary. The extra electron brought by proton will contribute to the conduction-band and raise the Fermi level, as shown in Figure 12c [59]. These protons will stay balanced in equilibrium under vacuum. In the case of MH, the hydrogen storage capability of a hydride former metal is related to its own electron density. For the study in MH from elements, the heat of hydride formation (indicator for M–H bond strength) and the work function for the first row of transition metals (from Sc to Ni) were plotted against their number of 3d electron in Figure 13. As the number of electrons increases, the work function increases due to the increase in the charge of the nucleus, and the host metal starts to resist incorporation of extra electrons brought by the absorbed hydrogen and consequently weakens the MH bond strength (less negative ∆H), with the exception of Mn. Mn has an extraordinarily low work function due to its containing of the maximum number of un-paired electrons (five) and also to a lower ∆H than predicted by the trend. For MH from intermetallic alloys, the situation is the opposite. Comparing the alloys LaNi₂ and LaNi₅, for example, the former has a lower electron density (lower Ni-content), a lower E_F (as M1 in Figure 13), and a tendency to trap the residual hydrogen near the interface. According to Equation (3), LaNi₂ (33% of La) has a stronger MH bond strength compared to that of LaNi₅ (16% of La). Therefore, we believe the phase with a stronger MH bond will keep the MH (β-phase) during a total degassing in vacuum. In this study, C14 and Ti₂Ni are deemed to be the phases with the strongest MH bond strength, which agrees with the prediction from Equation (3). These pre-occupied sites at the grain boundary on the side with a higher work function will act as the nucleation center for the β-phase growth (as illustrated in Figure 14). Therefore, the PCT absorption/desorption hysteresis of a single-phase AB₅ MH alloy is always larger than that from a typical multi-phase AB₂ MH alloy [7].

A sample was removed from the PCT apparatus with an equilibrium hydrogen pressure of 0.005 MPa, a pressure that separates the two plateaus in the PCT isotherms, as shown in Figure 5. From previous studies in BCC MH alloys, we believe the oxide layer formed at the alloy surface during the sudden exposure to air can prevent the loss of hydrogen into atmosphere [61]. Its XRD pattern is shown in Figure 11c, and the Rietveld analysis shows an unchanged BCC phase with 2/3 hydrided TiNi and Ti₂Ni phase (Table 8) and demonstrates that the higher plateau (from 0.1 MPa to 1 MPa) belongs to the BCC phase. Another sample was removed at the maximum hydrogen pressure (4 MPa). The XRD pattern of this sample shows a 90% transition from BCC (M) to FCC (MH) in the BCC phase and complete α- to β-phase transitions for the other three phases.

Besides the phase transition in the gaseous phase hydrogen reaction, we also performed XRD analysis on two pieces of electrode made from the same P17 alloy by induction melting with 50% and 100% state of charge (SOC). The resulting patterns are shown in Figure 11e,f. The Rietveld analysis shows a clear α-BCC phase and almost fully hydrided TiNi, Ti₂Ni, and C14 phases for the 50% SOC sample, while a 40% hydrided BCC (β-BCC) phase with the other phases fully hydrided are seen for the 100% SOC. During charging, the phases with stronger MH bond strength (more negative ∆H) were hydrided first, and BCC was charged last, as in the case of the gaseous phase reaction. The BCC phase was not hydrogenated before completion of hydride formation in other non-BCC phases. The BCC phase was never fully charged in the half-cell testing, due to the nature of the open-to-air configuration used in the electrochemical study. This explains the low electrochemical discharge capacity found in Figure 9, when compared to the gaseous phase maximum capacity.

There are some important differences in hydrogen storage between gaseous phase and electrochemistry, especially when the PCT isotherm (Figure 5) is compared with the electrochemical charge and discharge curves (Figure 15). While the gaseous hydrogen charging occurs through hydrogen gas molecule adsorption and splitting into two hydrogen atoms at the clean surface of metals (free from oxide), the electrochemical hydrogen charging is implemented by an applied voltage, which drives electrons into the negative electrode and leads to water splitting at the surface (which is typically covered with a thin oxide [62]). In PCT, hydrogen absorption can begin as soon as any amount of hydrogen gas is available. However, in the case of an electrochemical environment, a minimal voltage is required to generate enough of a strong electric field to split water into protons and hydroxide ions. Also in PCT, the equilibrium pressure depends only on the concentration of hydrogen in the MH; the electrochemical voltage is mainly determined by the surface reaction and thus a sudden change in the charge voltage profile is not seen when the active storing material switches from one phase into another one (at approximately 50% SOC). The discharge processes between the gaseous phase and electrochemistry are also different. For the gaseous phase, the movement of proton is influenced by diffusion and an equilibrium is reached when the same amount of hydrogen gas leaves and enters the metal. The electrochemical case is far more complicated. During discharge, the electron moves away from the MH alloy into the current collector, which forces the proton to move in the opposite direction, reaching the surface. Unless the surface recombination of either proton-hydroxide or proton-oxygen is very slow (which is unlikely) and there is a large number of protons accumulated at the electrolyte interface, the proton should continue to arrive at the interface to reach charge neutrality. Energy is gained from both the reduction in number of electrons in the metal system and the dehydride process. Therefore, there is no sudden drop in the discharge voltage profile when the system finishes discharging for one phase and starts the discharge for the next phase. This finding is very encouraging with regard to the electrochemical applications of MH alloys. In a multi-phase system, as long as there is an electrochemically active phase, the so called “irreversible” or “no participation” phases found in PCT can be fully utilized in the electrochemical environment. This finding can be used to explain the unpredicted high electrochemical capacity (compared to the gaseous phase capacities) observed in multi-phase systems involving the Zr₂Ni₇ phase [8,9].

With the knowledge that the visible plateaus of the PCT isotherm in the 0.1–1 MPa range originate entirely from the phase transition from BCC to FCC, the mechanism of this step, at equilibrium pressure in pure BCC alloys [61,63], is not well understood. Evidently, with synergetic effects through the non-BCC phases, the intermediate body-centered-tetragonal (BCT) phase is skipped. Although the hysteresis of the BCC phase is still very large (compared to all other MH alloy families), it has been significantly reduced through the synergetic effects of the non-BCC phases. In this study, it is clearly observed that, during hydrogenation, the non-BCC phases hydrogenate first and expand the lattice on the non-BCC phase side of the grain boundary. The lattice of the BCC phase is also pre-expanded before hydrogenation by the stress from the hydride on the other side, which reduces the energy barrier needed to expand the BCC lattice, and thus reduces the hysteresis and changes the FCC β-phase directly.

After we identified BCC as the storage phase and the others (C14, TiNi, and Ti₂Ni) as the catalytic phase, a correlation study between the properties of the constituent phases and the hydrogen storage characteristics (both in gaseous phase and electrochemistry) was conducted; the results are summarized in

Table 9. Table of correlation coefficient (R²) between hydrogen storage properties and phase component characteristics. Significant correlations are highlighted in bold.

Property	Gaseous Phase					Electrochemistry
	Plateau Pressure @30 °C	Plateau Pressure @60 °C	Maximum Capacity	Reversible Capacity	PCT Hysteresis	High-Rate Discharge Capacity	Full Discharge Capacity	HRD
BCC, lattice constant	0.02	0.04	0.14	0.36	0.05	0.07	0.04	0.06
BCC, crystallite size	0.00	0.04	0.21	0.26	0.40	0.57	0.15	0.53
BCC, abundance	0.11	0.01	0.22	0.41	0.53	0.38	0.27	0.28
C14, unit cell volume	0.17	0.08	0.01	0.72	0.17	0.09	0.02	0.08
C14, crystallite size	0.31	0.32	0.13	0.00	0.04	0.04	0.00	0.06
C14, abundance	0.10	0.08	0.05	0.09	0.07	0.19	0.02	0.26
TiNi, lattice constant	0.03	0.05	0.13	0.26	0.08	0.17	0.05	0.16
TiNi, crystallite size	0.16	0.15	0.09	0.19	0.06	0.04	0.02	0.04
TiNi, abundance	0.02	0.00	0.31	0.57	0.66	0.64	0.15	0.59

. Those with significant correlation (larger R²) are highlighted. There are two noticeable correlations in the gaseous phase properties. The first is the reversible capacity, which decreases with the increase in TiNi phase abundance and C14 unit cell volume (Figure 16a). The former has already been explained by the TiNi phase being the largest component in the lower plateau (not seen from the PCT isotherm in Figure 5) and makes the largest contribution to the irreversible part of the gaseous phase hydrogen storage capacity (Figure 6). The source of the latter correlation is less clear. C14, although it has a relatively low ∆H, may be still a catalytic phase. With the decrease in the unit cell volume, hydride from the C14 phase becomes less stable and contributes positively to the reversible capacity of the storage phase (TiNi and BCC). The second noticeable correlation found in the gaseous phase properties is related to the PCT hysteresis—more precisely, the hysteresis of the BCC phase. It has been found that the 30 °C PCT hysteresis increases with increasing TiNi phase abundance and decreasing BCC phase abundance (Figure 16b). Many studies have been conducted on the PCT hysteresis (for a review see [64]). It is generally accepted that the PCT hysteresis is the elastic energy needed for the deformation of the lattice near the α-β boundary. Higher PCT hysteresis in an MH alloy increases the difficulty of hydrogenation. From this aspect, the BCC phase facilitates reaction, while TiNi phase retards the hydrogenation of the BCC phase, which is very difficult to observe from the large value of hysteresis since it involves phase changes from BCC to FCC [6].

As for correlation to the electrochemical properties, both the BCC crystallite size and the TiNi phase abundance show significant correlation to the high-rate discharge performance. A smaller crystallite in the BCC phase is more desirable, as it gives better high-rate performance (Figure 16c). It is possible that smaller crystallite sizes will generate more grain boundaries (shown in Figure 12c) and facilitate more charge transfer between the two phases. The phase abundances from the three major phases are plotted against HRD in Figure 16d. The general trends suggest that the BCC and C14 phases are more beneficial to HRD, while the TiNi phase deteriorates it.

5. Conclusions

In this study on the synergetic effects in MH alloys, we challenge the argument that the catalytic phase must have a weaker MH bond strength compared to that in the main storage phase. The C14 phase, with an ostensibly stronger MH bond, contributes to the reversibility in the hydrogen storage of other storage phases. The BCC phase, in a multi-phase system, demonstrates a relatively weak MH bond. On the other hand, the TiNi phase, also with a strong MH bond, hinders the hydrogenation of the BCC phase. The driving force for proton movement in the alloy varies for gaseous phase and electrochemistry. The former occurs by diffusion and the equilibrium with gaseous hydrogen at the clean surface, while the latter proceeds via the electric field and the local distribution of valence electrons. The synergetic effects in a multi-phase MH system can be explained by the pre-occupied hydrogen sites on the side of the metal with a larger work function near the grain boundary.