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

The electrochemical reactions of multi-phase metal hydride (MH) alloys were studied using a series of Laves phase-related body-centered-cubic (BCC) Ti15.6Zr2.1V43Cr11.2Mn6.9Co1.4Ni18.5Al0.3X (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.


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.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, Zr 7 Ni 10 , 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.In, TiMn 2 additives Desorption kinetics GP [30] * indicates that the authors attributed improvement to micro-cracking at the surface instead of to synergetic effects.
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 Ti 2 Ni, 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 15.6 Zr 2.1 V 44 Cr 11.2 Mn 6.9 Co 1.4 Ni 18.5 Al 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.

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.

X-Ray Diffraction Structure Analysis
Eight alloys with the targeted compositions listed in Table 2 were prepared by arc melting.Th 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 AB 2 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.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  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 AB 2 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).more popular (larger abundance), its crystallite becomes larger.In the case of BCC and C14 phases, the trend is the opposite (Figure 3).

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.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 (hyperstoichiometry), 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 hyperstoichiometry.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  more popular (larger abundance), its crystallite becomes larger.In the case of BCC and C14 phases, the trend is the opposite (Figure 3).

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.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 (hyperstoichiometry), 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 hyperstoichiometry.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

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.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 Ti 2 Ni 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 (Ti 2 Ni).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.

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 AB 2 [46], AB 5 [47], and A 2 B 7 [27] MH alloy families.The gaseous phase properties obtained from the PCT analysis are summarized in Table 5.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.
Batteries 2016, 2, 15 8 of 21 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 (∆H) and entropy (∆S), were calculated from the equation for free energy (G): 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 ∆H and ∆S were calculated and listed in Table 6.Conventionally, the values obtained using the desorption isotherm (∆H D and ∆S 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 ∆S 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, ∆H D is always lower than ∆H A , and ∆H (A+D)/2 falls between the two values.∆S D can be higher or lower than ∆S A , and ∆S (A+D)/2 also falls in the middle.
In our previous publication, we adopted an estimation method for ∆H from the atomic percentage and obtained a ∆H of the three hydride-former atoms, Zr (∆H h = ´163 kJ¨mol ´1 H 2 for ZrH 2 ), Ti (∆H h = ´124 kJ¨mol ´1 H 2 for TiH 2 ), and V (∆H h = ´34 kJ¨mol ´1 H 2 for VH 2 ).This estimation is based on the following assumptions.The complete heat of hydride formation of AB n alloy should be: However, the ∆H for the alloy AB n and the non-hydride former atoms are smaller than those in the hydride former atoms (for example, ∆H for VCr 2 and NiH 0.5 are ´2 kJ¨mol ´1 H 2 [51] and ´6 kJ¨mol ´1 H 2 [52], respectively).Therefore, the heat of hydride formation can be estimated by: 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 ∆H for the hydride of the constituent phases of alloy P17 are ´28.5 (BCC), ´55.4 (TiNi) and ´49.3 kJ¨mol ´1 H 2 (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.
Batteries 2016, 2, 15 9 of 21 Thermodynamic properties, including changes in enthalpy (ΔH) and entropy (ΔS), were calculated from the equation for free energy (G): 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.Conventionally, the values obtained using the desorption isotherm (ΔHD and ΔSD) 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 ΔSA 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, ΔHD is always lower than ΔHA, and ΔH(A+D)/2 falls between the two values.ΔSD can be higher or lower than ΔSA, 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 (∆Hh = −163 kJ•mol −1 H2 for ZrH2), Ti (∆Hh = −124 kJ•mol −1 H2 for TiH2), and V (∆Hh = −34 kJ•mol −1 H2 for VH2).This estimation is based on the following assumptions.The complete heat of hydride formation of ABn alloy should be: However, the ΔH for the alloy ABn and the non-hydride former atoms are smaller than those in the hydride former atoms (for example, ΔH for VCr2 and NiH0.5 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: 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.The ∆Hs 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 ∆H values in Figure 7a.Except for alloy P42 (Mo), the measured ∆Hs, 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 ∆H is due to the synergetic effects associated with the other two phases that have a much lower ∆H 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 ∆H 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) ∆H, 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 ∆H, Mo has a positive contribution (∆H for MoH 0.5 = +10 kJ¨mol ´1 H 2 [52]).This explains the abnormal behavior of alloy P42 (Mo) in Figure 7.  4) using Equation (3).

Electrochemical Properties
The electrochemical hydrogen storage properties were determined in a flooded half-cell configuration with a commercially available co-precipitated Ni91Co4.5Zn4.5 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 cutoff 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.It is obvious that while the increase  4) and Equation (3).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).

Electrochemical Properties
The electrochemical hydrogen storage properties were determined in a flooded half-cell configuration with a commercially available co-precipitated Ni 91 Co 4.5 Zn 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.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% H 2 = 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).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).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 In order to further study HRD performance, the bulk diffusion constant (D) and surface exchange current (I 0 ) 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 I 0 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 I 0 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 pvs.NiOOH at 0.36 Vq " 1.294 `0.029logppH 2 q r7s (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].
Batteries 2016, 2, 15 12 of 21 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].

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)

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 (Zr 7 Ni 10 , for example [57]), and the secondary phase with a higher plateau pressure and less negative ∆H is considered to be the catalytic phase (Zr 2 Ni 7 , 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 Ti 2 Ni 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% Ti 2 Ni, and 6.2% C14 (Table 8).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 Ti 2 Ni.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 LaNi 2 and LaNi 5 , 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), LaNi 2 (33% of La) has a stronger MH bond strength compared to that of LaNi 5 (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 Ti 2 Ni 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 AB 5 MH alloy is always larger than that from a typical multi-phase AB 2 MH alloy [7].
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 Figures 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) Figure 13.A plot of the heat of hydride formation (data from [52]) and work function (data from [60]) vs. the number of 3d electrons for the first row of transition metal elements from Sc to Ni.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 In this illustration, the M1 phase has a larger W than M2, and contains a small amount of hydride near the boundary, even at a very low SOC.The MH (β) phase nucleates from a randomly distributed proton (α) in the gaseous phase (left) that, during electrochemistry, starts from the remaining MH in the phase with a larger W near the grain boundary.
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 Ti 2 Ni 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 αto 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, Ti 2 Ni, 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 Zr 2 Ni 7 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 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 Ti 2 Ni) 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.Those with significant correlation (larger R 2 ) 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.

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.

Figure 2 .
Figure 2. Plots of the lattice constants a for the TiNi, BCC, and C14 phases vs. (a) the metallic radius of the additives and (b) the Mn-content in the BCC phase.

Figure 2 .
Figure 2. Plots of the lattice constants a for the TiNi, BCC, and C14 phases vs. (a) the metallic radius of the additives and (b) the Mn-content in the BCC phase.

Figure 2 .
Figure 2. Plots of the lattice constants a for the TiNi, BCC, and C14 phases vs. (a) the metallic radius of the additives and (b) the Mn-content in the BCC phase.
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 Figure2b.Other correlations are less obvious.
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 Figure5should 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).

Figure 6 .
Figure 6.Plot of the amount of irreversible hydrogen storage found in PCT analysis vs. the TiNi phase abundance.

Figure 6 .
Figure 6.Plot of the amount of irreversible hydrogen storage found in PCT analysis vs. the TiNi phase abundance.

Figure 7 .Table 6 .
Figure 7. (a) The measured ∆H based on the absorption, half-point, and desorption isotherms and (b) Plateau pressure and BCC phase lattice constants vs. the calculated ∆H based on BCC composition determined by EDS (Table4) and Equation (3).

Figure 7 .
Figure 7. (a) The measured ∆H based on the absorption, half-point, and desorption isotherms and (b) Plateau pressure and BCC phase lattice constants vs. the calculated ∆H based on BCC composition determined by EDS (Table4) and Equation (3).

Figure 8 .
Figure 8.The evolution of (a) maximum capacity and (b) HRD in the first 13 cycles.

Figure 9 .
Figure 9.Comparison of capacities obtained from gaseous phase and electrochemical measurements.The gaseous phase capacity is converted by 1 wt% = 268 mAh•g -1 .

Figure 8 .
Figure 8.The evolution of (a) maximum capacity and (b) HRD in the first 13 cycles.

Figure 8 .
Figure 8.The evolution of (a) maximum capacity and (b) HRD in the first 13 cycles.

Figure 9 .
Figure 9.Comparison of capacities obtained from gaseous phase and electrochemical measurements.The gaseous phase capacity is converted by 1 wt% = 268 mAh•g -1 .

Figure 9 .
Figure 9.Comparison of capacities obtained from gaseous phase and electrochemical measurements.The gaseous phase capacity is converted by 1 wt% = 268 mAh¨g ´1.

Figure 10 .
Figure 10.The plots of gaseous phase maximum hydrogen storage capacity and equilibrium plateau pressure from PCT desorption isotherm vs. the OCV obtained from the half-cell measurement.

Figure 10 .
Figure 10.The plots of gaseous phase maximum hydrogen storage capacity and equilibrium plateau pressure from PCT desorption isotherm vs. the OCV obtained from the half-cell measurement.

Figure 11 .
Figure 11.XRD patterns of a P17 alloy prepared by the conventional melt-and-cast process: (a) pristine; (b) hydrided at 4 MPa H2 gas and then degassed in vacuum for 8 h at 300 °C; (c) withdrawn from an equilibrium state with 0.005 MPa H2 pressure; and (d) pull-out from an equilibrium state with 4 MPa H2 pressure, and the electrodes made from the same alloy with (e) 50% and (f) 100% state of charge achieved electrochemically.

Figure 11 .
Figure 11.XRD patterns of a P17 alloy prepared by the conventional melt-and-cast process: (a) pristine; (b) hydrided at 4 MPa H 2 gas and then degassed in vacuum for 8 h at 300 ˝C; (c) withdrawn from an equilibrium state with 0.005 MPa H 2 pressure; and (d) pull-out from an equilibrium state with 4 MPa H 2 pressure, and the electrodes made from the same alloy with (e) 50% and (f) 100% state of charge achieved electrochemically.

Figure 12 .
Figure 12.Two MH alloys with different work functions (W): (a) before making contact; (b) after making contact; and (c) after hydrogenation and dehydrogenation.

Figure 12 .
Figure 12.Two MH alloys with different work functions (W): (a) before making contact; (b) after making contact; and (c) after hydrogenation and dehydrogenation.

Figure 12 .
Figure 12.Two MH alloys with different work functions (W): (a) before making contact; (b) after making contact; and (c) after hydrogenation and dehydrogenation.

Figure 13 .
Figure13.A plot of the heat of hydride formation (data from[52]) and work function (data from[60]) vs. the number of 3d electrons for the first row of transition metal elements from Sc to Ni.

Figure 14 .
Figure 14.Comparison of hydrogen storage through (a) a gaseous phase reaction in a perfect singlecomponent system and (b) an electrochemical reaction in a two-phase system with different work functions (WM1 > WM2) at different states of charge (SOC).In this illustration, the M1 phase has a larger W than M2, and contains a small amount of hydride near the boundary, even at a very low SOC.The MH (β) phase nucleates from a randomly distributed proton (α) in the gaseous phase (left) that, during electrochemistry, starts from the remaining MH in the phase with a larger W near the grain boundary.

Figure 14 .
Figure 14.Comparison of hydrogen storage through (a) a gaseous phase reaction in a perfect single-component system and (b) an electrochemical reaction in a two-phase system with different work functions (W M1 > W M2 ) at different states of charge (SOC).In this illustration, the M1 phase has a larger W than M2, and contains a small amount of hydride near the boundary, even at a very low SOC.The MH (β) phase nucleates from a randomly distributed proton (α) in the gaseous phase (left) that, during electrochemistry, starts from the remaining MH in the phase with a larger W near the grain boundary.
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].

Figure 15 .
Figure 15.The second cycle electrochemical charge and discharge voltage profiles for an electrode made from alloy P40.The smooth curves seen here are very different from the step-shape PCT isotherm shown in Figure 5.

Figure 15 .
Figure 15.The second cycle electrochemical charge and discharge voltage profiles for an electrode made from alloy P40.The smooth curves seen here are very different from the step-shape PCT isotherm shown in Figure 5.

Figure 16 .
Figure 16.Plots of (a) C14 unit cell volume and TiNi phase abundance vs. gaseous phase reversible capacity; (b) TiNi and BCC phase abundance vs. PCT hysteresis measured at 30 ˝C; (c) HRD vs. BCC phase crystallite size; and (d) TiNi, BCC, and C14 phase abundances vs. HRD.

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.

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

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

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

Table 4 .
Energy-dispersive spectroscopy (EDS) composition data from select spots in Figure4.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.

Table 5 .
Summary of gaseous phase and thermodynamic properties.

Table 5 .
Summary of gaseous phase and thermodynamic properties.

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.

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.

Table 7 .
Summary of electrochemical properties.Cap 100 and Cap 8 are discharge capacities measured for discharge currents at 100 mA¨g ´1 and 8 mA¨g ´1, respectively.HRD 100 is the ratio of capacities measured at 100 mA¨g ´1 and 8 mA¨g ´1.RT: room temperature; and OCV: open circuit voltage.

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.

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.

Table 9 .
Table of correlation coefficient (R 2 ) between hydrogen storage properties and phase component characteristics.Significant correlations are highlighted in bold