Influence of Pd on the structure and electrochemical hydrogen storage properties of Mg50Ti50 alloy prepared by ball milling

Electrochimica Acta 55 (2010) 611–619

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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta

Influence of Pd on the structure and electrochemical hydrogen storage properties of Mg50 Ti50 alloy prepared by ball milling Steeve Rousselot, Anais Gazeau, Daniel Guay, Lionel Roué ∗ INRS-Énergie, Matériaux et Télécommunications, 1650 Blvd. Lionel-Boulet, Varennes (Québec), Canada J3X 1S2

a r t i c l e

i n f o

Article history: Received 8 July 2009 Received in revised form 31 August 2009 Accepted 6 September 2009 Available online 16 September 2009 Keywords: Metal hydride Mg–Ti alloys Palladium Mechanical alloying Ni-MH batteries

a b s t r a c t High-energy ball milling was used to modify the physico-chemical and the electrochemical hydrogenation properties of Mg50 Ti50 alloy via the addition of Pd. This was done by first ball milling Mg and Ti together for (20 − x) hours. 3.3 at.% Pd was then added and ball milling was resumed for x hours. X-ray diffraction and X-ray photoelectron spectroscopy analyses revealed that the alloying of Pd with pre-milled Mg50 Ti50 was initiated after only a few minutes and was completed after 5 h of milling. The maximum discharge capacity of the Mg50 Ti50 –3.3 at.% Pd electrode increased significantly with the milling time (from 35 mAh g−1 for 5 min to 480 mAh g−1 for 20 h of milling). The exchange current density increased with the milling time and was directly related to the Pd surface concentration, suggesting that Pd plays a key role in facilitating the charge-transfer reaction. In contrast, the incorporation of Pd had a minor effect on the hydrogen diffusion coefficient. The electrochemical pressure-composition isotherms revealed a significant destabilization of the hydride as the milling time with Pd increased. No significant improvement in the hydrogen storage properties of Mg50 Ti50 –Pd electrodes was observed for Pd concentrations higher than 3.3 at.%. © 2009 Elsevier Ltd. All rights reserved.

1. Introduction Much attention has been recently paid to the Mg–Ti–H system from both fundamental and practical viewpoints [1–21]. Indeed, due to their high hydrogen diffusivity and solubility (up to 6.5 wt.% H for Mg80 Ti20 ) [1], Mg–Ti alloys are promising candidates as negative electrode materials for Ni-MH batteries [1–11] and as hydrogen storage materials [12–15]. Moreover, their remarkable H-induced optical properties offer new possibilities as switchable mirrors for smart solar collectors [16–18] and as optical hydrogen detectors [19]. The face-centered cubic (fcc) structure of the hydride phase of Mgx Ti(1−x) alloys with x ≤ 85 at.% is most likely responsible for their excellent hydrogen storage properties [20,21]. The synthesis of Mg–Ti alloys is a major challenge considering that the solid solubility of Ti in Mg is close to zero and that there are no intermetallic compounds, as illustrated by the Mg–Ti phase diagram [22]. Moreover, conventional metallurgical methods are hardly applicable because the melting point of Ti (1943 K) significantly exceeds the boiling point of Mg (1363 K). Metastable single-phase Mg–Ti thin films have been successfully synthesized over a large stoichiometric range by means of electron-beam and magnetron co-sputter deposition techniques [1–4,14–21]. How-

∗ Corresponding author. Tel.: +1 450 929 8185; fax: +1 450 929 8102. E-mail address: [email protected] (L. Roué). 0013-4686/$ – see front matter © 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2009.09.014

ever, these synthesis methods are not adapted for the large scale preparation of low cost bulk Mg–Ti materials as required for NiMH battery and fuel cell applications. In this context, mechanical alloying has demonstrated its high efficiency for the room temperature elaboration of a wide range of non-equilibrium Mg–Ti materials starting from elemental or hydrided Mg and Ti powders [5–13]. When prepared by mechanical alloying, Mg–Ti materials must be activated for electrochemical hydrogen storage by adding of few at.% of Pd [5–9]. Indeed, in the absence of Pd, the hydrogen discharge capacity of ball-milled Mg–Ti materials is close to zero [7]. This is a major disadvantage considering the high cost of Pd and thus, studies must be undertaken to minimize its content or to replace it by a less expensive hydrogenation catalyst. As a first step toward this goal, the role of Pd in the electrochemical H-sorption of ballmilled Mg–Ti materials as well as its incidence on the Mg–Ti alloy structure need to be clarified. This is the aim of the present work. For this reason, Mg50 Ti50 –Pd materials were prepared by ball milling with 3.3 at.% of Pd added after various periods of milling. Their crystalline structure and chemical surface state were investigated by X-ray diffraction (XRD) and X-ray photoelectron spectroscopy (XPS). Their electrochemical hydrogen sorption properties were characterized in terms of hydrogen discharge capacities, charge-transfer, H-diffusion kinetics and H-absorption thermodynamic properties. Finally, attempts to replace Pd by other hydrogenation catalysts will be discussed.

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2. Experimental

Pure Mg (99.9%, chips, Norsk Hydro) and Ti powders (99.5%, −325 mesh, Alfa Aesar) were introduced with a Mg:Ti atomic ratio of 50:50 into a cylindrical stainless steel crucible (55 ml) with three stainless steel milling balls (one of 14 mm and two 11 mm diameter), corresponding to a ball-to-powder mass ratio of 10:1. Pd powder (3.3 at.% corresponding to 9 wt.%, 99.9%, −325 mesh, Alfa Aesar) was added following two procedures. In the first one, Mg and Ti were first milled for 20 − x hours before the addition of Pd. The milling operation was then resumed for x hours in presence of Pd, so that the total milling time was fixed at 20 h. In the second procedure, Pd was introduced simultaneously with the Mg and Ti powders at the beginning of the milling and the milling time was set to 20 h. Additionally, some samples were prepared by varying the Pd content (from 0 to 6.6 at.%) and by replacing Pd by other elements (Pt, Ni, Nb, V), using the second procedure. In all cases, the crucible was sealed under an Ar atmosphere. The powders were milled using a vibratory type mill (Spex 8000 M). The milling yields (defined as the ratio of the powder masses after and before milling) were higher than 70% in all cases, indicating the absence of excessive cold welding between the powder particles and the milling tools. Energy dispersive X-ray (EDX) analysis of the milled powders showed that the Fe content was less than 1 at.%, reflecting the limited erosion of the container and balls. On the basis of scanning electron microscopy micrographs (not shown), the morphologies of the different milled powders are similar and consist of irregular particles, with a grain size distribution ranging from less than 1 ␮m to several tens of ␮m.

trode. The working electrode was charged at a current density of −200 mA g−1 for 3 h and discharged at 20 mA g−1 followed by a deeper discharge at 5 mA g−1 up to −0.4 V vs. Hg/HgO. The discharge capacities are reported in mAh g−1 of active material (i.e., the masses of Mg, Ti and Pd were considered for the calculation of the discharge capacity). Linear polarization and potential step experiments on the composite electrodes were measured on a Voltalab40 (Radiometer Analytical) potentiostat/galvanostat/FRA apparatus. Prior to the linear polarization measurements, the electrode was fully charged and rested for 2 h up to stabilization of the open circuit potential. Linear polarization was performed by scanning the electrode potential from −10 to 10 mV (vs. open circuit potential) at 0.1 mV s−1 . Curves were corrected for the ohmic drop determined from impedance measurements. The potential step experiments were carried out on electrodes whose active material particles were selected between 20 and 75 ␮m in size by sieving the milled powder. The electrode was first fully charged and rested for 2 h up to stabilization of the open circuit potential, then a potential step of 50 mV vs. open circuit potential for 45 min was applied and the discharge current was monitored as a function of time. Electrochemical hydrogen pressure-composition isotherms were performed at room temperature (23 ◦ C) on an Arbin BT2000 from equilibrium potential measurements in the discharge step. The electrode was first charged at −200 mA g−1 for 3 h. Then a discharge current of 5 mA g−1 was applied for 1 or 2 h depending on the discharge capacity of the material and the stabilized equilibrium potential was measured after an open circuit period of 2 h. The procedure was repeated until the electrode was completely discharged.

2.2. Material characterization

3. Results and discussion

The materials were characterized by X-ray diffraction (XRD) using a Bruker D8 diffractometer with Cu K␣ radiation. Structure refinement was performed according to the Rietveld method [23] using GSAS and EXPGUI softwares [24,25]. A total of three phases (hcp, bcc and fcc) with nine independent variables (four lattice parameters, three crystallite sizes and two phase proportions) were varied. The isotropic temperature factors (uiso value) of each phase were kept constant (uisohcp = 0.025, uisobcc = 0.18 and uisofcc = 0.025). This strategy is based on two previous works where a detailed description of the Rietveld refinement procedure is given and where the self-consistency of the results obtained is stressed [8,9]. The chemical surface state of the powder particles was analyzed by X-ray photoelectron spectroscopy (XPS) on a VG Escalab 2201XL spectrometer with an Al K␣ radiation monochromatic source. Surface contamination was removed by sputtering Ar+ ions of 3 keV energy. The C 1s core level peak of adventitious C contamination (284.6 eV) was used as internal reference. The Mg, Ti and Pd surface atomic concentrations were estimated from the Mg 1s, Ti 2p and Pd 3d core level peaks, using sensitivity factors of 11.18, 7.91 and 9.48, respectively. The surface concentrations were corrected considering the sputtering yield of each element under argon bombardment (8.29, 3.80 and 10.90% for Mg, Ti and Pd, respectively).

3.1. Structural analysis

2.1. Material synthesis

2.3. Electrochemical experiments Electrochemical charge/discharge cycling tests were carried out on an Arbin BT2000 battery tester at room temperature in a 6 M KOH electrolyte using a three-electrode cell. The working electrode was made of a mixture of 100 mg of active material, 800 mg of graphite and 20 mg of carbon black. The counter electrode was a nickel wire and the reference electrode was an Hg/HgO elec-

Fig. 1 depicts the X-ray diffraction patterns of Mg50 Ti50 powder milled with 3.3 at.% of palladium for different milling times (from 15 min to 20 h). In all cases, the characteristic diffraction peaks of the starting Mg or Ti powders are not found in the XRD patterns, confirming that alloying between Mg and Ti is complete. Instead, a relatively intense and broad peak is observed in the 2 region extending from 30◦ to 45◦ , with a series of weak peaks at higher 2 angle values. As far as we can tell from a visual inspection of Fig. 1, the broad XRD structure is not affected by the milling time in presence of Pd. Also observed in the XRD patterns of Fig. 1 are the characteristic diffraction peaks of Pd. These peaks are located at 2 = 40.1◦ , 46.6◦ , 68.1◦ , 82.1◦ and 86.6◦ , and they correspond to the (1 1 1), (2 0 0), (2 2 0), (3 1 1) and (2 2 2) reflections of pure fcc palladium, respectively. The intensity of these peaks decreases with milling. After 5 h of milling in presence of Mg50 Ti50 , the (1 1 1) reflection of Pd is barely visible while the others peaks are not detectable. For 20 h of milling, the diffraction peaks of Pd are absent from the XRD trace. A Rietveld refinement analysis was performed on the XRD patterns to determine the phase composition and phase proportion of the compounds. A detailed analysis of the different phases that are formed upon extensive milling of Mg and Ti has been reported previously [9], and the main results of this study is reported in the first entry of Table 1 (w/o Pd). Also, a detailed review of the literature data related to the formation of alloyed compounds between Mg and Ti upon milling was undertaken elsewhere [9]. So based on this information, three phases were used to fit the experimental data, namely a hexagonal-close-packed (hcp), a body-centeredcubic (bcc) and a face-centered-cubic (fcc) phase. The structural parameters extracted from the analysis of the XRD patterns are

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Table 1 Results of the Rietveld refinement analysis of Mg50 Ti50 –3.3 at.% Pd samples milled with Pd for different times. The total milling time is 20 h. Milling time with Pd

hcp

bcc

Constants

w/o Pd 15 min 30 min 1h 2h 5h 20 h

a (Å)

c (Å)

3.05 3.00 3.01 2.99 2.99 2.98 3.00

4.93 4.83 4.85 4.87 4.85 4.84 4.86

Volume (Å3 )

Crystal. size (nm)

wt.%

39.7 37.5 38.0 37.8 37.5 37.3 37.9

3 5 5 4 5 3 4

24 19 23 17 20 25 26

Constant

fcc (Pd)

Volume (Å3 )

Crystal. size (nm)

wt.%

41.0 39.5 39.2 39.5 39.1 39.1 40.1

4 3 3 3 3 4 5

76 78 74 80 78 74 74

a (Å) 3.45 3.41 3.40 3.40 3.39 3.39 3.42

summarized in Table 1, and the contribution of each phase to the experimental XRD pattern is depicted in Fig. 1. As seen in Table 1, there is no systematic variation of the lattice constants and phase proportion of the hcp and bcc phases with the milling time. Considering that the diffraction peaks are broad and not very intense, it can be concluded that the volume of the hcp and bcc phases remain constant at ∼37.5 and 39.5 Å3 , respectively, when the milling time in presence of Pd is increased. As it will be shown later on, the amount of crystalline Pd decreases with time as the milling operation is pursued. However, this gradual decrease of the fcc Pd content is not accompanied by a systematic change in

Constant

Volume (Å3 )

Crystal. size (nm)

wt.%

– 58.9 59.0 59.0 59.0 59.0 N/A

– 46 38 34 24 24 N/A

– 3 3 3 2 <1 N/A

a (Å) – 3.89 3.89 3.89 3.89 3.89 N/A

the lattice parameters of the hcp and bcc phases. It is not surprising considering the small number of Pd atoms that would be inserted in these phases (maximum of 3.3 at.% for 20 h of milling). As it will be seen later on, evidence of Pd alloying with Mg and Ti will be reached through a detailed analysis of X-ray photoelectron spectroscopy data. As reported elsewhere for material with similar composition [9], the hcp phase is a Ti-rich phase, with Ti occupying more than 80% of sites of the unit cell. In contrast, the lattice parameter of the bcc phase suggests that it contains an almost equal proportion of Mg and Ti atoms. The crystallite sizes of both hcp and bcc phases are constant (4 ± 1 nm). In contrast, there is a systematic variation of the crystallite size of Pd with the milling time: it decreases from 46 nm after 15 min of milling to 24 nm after 5 h. The lattice parameter of fcc Pd does not vary with the milling time and remains constant at 3.89 Å, which corresponds to the value expected for pure Pd. After 15 min of milling, there is only 2–3 wt.% Pd phase, and this value stays constant up to 5 h of milling, where it decreases to <1 wt.%. After 15 min of milling, the Pd content is curiously low compared to the nominal value (9 wt.%). The discrepancy between both values might originate from the fact that alloying of Pd with either or both of the hcp and bcc phase occurs on a very short time scale. As we will see later on, this hypothesis is consistent with the very rapid change observed in the XPS spectra after 15 min of milling. Pure Pd is not detected anymore after 20 h of milling, suggesting its complete dissolution in the Mg50 Ti50 alloy. This is in accordance with the fact that all of our efforts to identify free fcc Pd in 20 h-milled materials through transmission electron microscopy analysis failed. 3.2. Surface composition analysis

Fig. 1. X-ray diffraction patterns and Rietveld refinement of milled Mg50 Ti50 –3.3 at.% Pd. All samples were pre-milled for (20 − x) hours before the introduction of Pd, and the milling operation was then resumed for x hours. The milling time with Pd is indicated on each curve.

The surface composition of the samples was analyzed by Xray photoelectron spectroscopy (XPS). Carbon contamination at the surface of the samples was removed by Ar sputtering (20 min). In each case, iron was below the detection limit, indicating that contamination by the milling tools is minimal. Fig. 2A displays the Pd 3d5/2 and 3d3/2 core level regions of 20 h-milled Mg50 Ti50 –3.3 at.% Pd samples as a function of milling time with Pd. For short milling durations, both core level regions exhibit a doublet, indicating that Pd atoms exist in two different chemical states at the surface of the samples. These doublets were fitted using two components with the same Full Width at Half Maximum (FWHM = 1.1 eV). For the Pd 3d5/2 core level peak, these two components, Pd(A) and Pd(B), are located at ca. 335.2 and 336.1 eV, while they occur at 340.5 and 341.4 eV for the Pd 3d3/2 core level peak, respectively. The energy difference between each doublet is ∼5.3 eV, as expected for Pd. As seen in Fig. 2A, the relative importance of both components varies with the milling time and Fig. 2B shows how the ratio Pd(B)/Pd(A) + Pd(B) varies with the milling time. After 5 min

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Fig. 2. (A) X-ray photoelectron spectroscopy 3d core level peaks of Mg50 Ti50 –3.3 at.% Pd. The milling time in presence of Pd is indicated on each curve. (B) Variation of the relative content Pd(B)/Pd(A) + Pd(B) with the milling time in presence of Pd. (C) Variation of the total Pd surface content with the milling time in presence of Pd.

of milling, Pd(B) accounts for ∼40% of the surface Pd atom, and this proportion increases rapidly to 100% after 5 h of milling. In contrast, the Pd 3d5/2 and 3d3/2 core level peaks of pure Pd milled in the same conditions displays only a single component centered at 335.1 and 340.5 eV, respectively (not shown). The position of the Pd 3d5/2 and 3d3/2 core level peaks does not change with time even if the milling operation is pursued during 5 h. This component corresponds to Pd(A) described above. This means that component (A) of Fig. 2A corresponds to un-alloyed Pd, and that component (B) belongs to a different chemical form of Pd whose relative importance with respect to pure Pd increases gradually with the milling time. Component (B) of both Pd 3d5/2 and 3d3/2 core level peaks is shifted by 1.0 eV toward higher binding energies compared to component (A). Three factors could explain this shift: (i) crystallite size effect, (ii) formation of an oxide and (iii) alloying effect. It has been reported that the binding energy of pure metal cluster decreases as the cluster size decreases to the nanometer range [26–29]. The same phenomenon has been observed for amorphous Pd [30,31] or when the average coordination number of surface Pd atoms decreases with respect to metallic Pd [27]. In the case of isolated Pt cluster at the surface of either highly oriented pyrolytic graphite (HOPG) or gold substrate, this effect is hardly seen for cluster with a diameter larger than 3 nm, and a ∼0.6 eV binding energy difference is observed for isolated clusters with diameter less than 2 nm [32]. As shown in Table 1, the crystallite size of fcc Pd is of the order of 20–40 nm, much larger than the value quoted above. Also,

component (B) is not observed in the Pd 3d5/2 and 3d3/2 core level peaks of milled Pd alone (not shown), ruling out the possibility that crystallite size reduction could be responsible for the apparition of component (B). The formation of a palladium oxide at the surface of the material could also explain the apparition of a component at higher binding energy. The O 1s core level peak occurs in the same energy range as the Pd 3p3/2 peak, and it is difficult to differentiate both species. However, the characteristic KLL Auger peaks of O atoms at 978, 999 and 1013 eV are clearly discernible on the survey scan, indicating that contamination from O atoms are still present. While oxidation of Pd during milling cannot be totally ruled out, it is expected that oxidation of Mg or Ti will occur preferentially during milling as these two elements are much more sensitive to oxygen than Pd. This has been nicely demonstrated when RuO2 and Ti are milled together. In this case, an oxydo-reduction reaction is taking place giving rise to the formation of metallic Ru and oxidized Ti [33]. Also, as stated above, component (B) is not observed in the Pd 3d5/2 and 3d3/2 core level peaks of milled Pd, ruling out the possibility that oxidation of Pd could occur during milling. It is hypothesized that the apparition of component (B) in the Pd 3d5/2 and 3d3/2 core level peaks of Mg50 Ti50 –3.3 at.% Pd reflects the fact that Pd is alloyed with Mg and Ti. Reports in the literature indicating a positive (higher binding energy) shift of the core level peaks of Pd upon alloying with other elements abound. For example, a positive shift ranging from 0.8 to 1.5 eV has already been observed for Pd alloyed with other metals such as Ti [27], Mg [34],

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Sn [35] or Cu [36]. On this basis, component (B) in the Pd 3d5/2 and 3d3/2 core level peaks of Mg50 Ti50 –3.3 at.% Pd could then reasonably originate from the incorporation (alloying effect) of Pd in either or both the hcp and bcc phases that composed the Mg50 Ti50 matrix. Component (B) of the Pd5/2 core level peak shows a peak shift toward lower binding energy with increase milling time. The reason underlying this small change is not known for the moment. The fact that the relative proportion of Pd(B) is so large even at short milling time (see Fig. 2B) would indicate that the alloying reaction occurs on a very short time scale. As mentioned previously, the fact that Mg and Ti are pre-milled could be responsible for this effect. The high negative enthalpies of mixing between Pd and Mg or Ti [37,38] are also favourable conditions for the alloying reaction. As far as we can tell, the growth of component (B) is not associated with any modification in the Ti 2p and Mg 1s core level spectra. However, considering the very low Pd content (3.3 at.%), it seems reasonable that the alloying effect is more easily observed in the less concentrated element. The Pd surface concentration of Mg50 Ti50 –3.3 at.% Pd samples was calculated from the XPS data by taking into account the Mg 1s, Ti 2p3/2 and Pd 3d5/2 core level peaks. Thus, the relative Pd surface content is plotted in Fig. 2C with respect to the milling time with Pd. The Pd surface concentration increases from ∼0.5 at.% at short milling time to ∼3.3 at.% after 20 h of milling, which corresponds to the nominal concentration of Pd in the sample. The fact that only 0.5 at.% of Pd is detected at the surface of the sample at short milling time may be related to the difference in the ductility of Mg50 Ti50 and Pd. As explained by Suryanarayana [39], in the early stage of milling, brittle particles tend to be occluded and trapped by the more ductile particles. This would lead to the depletion of Pd atoms from the surface of the particles at short milling duration. With further milling, this phenomenon becomes less important as the brittle particles get uniformly distributed among the ductile matrix or chemical homogeneity is achieved between both components. In both cases, the chemical composition of the material approaches the nominal composition. This is exactly what is observed in this case. Ma et al. [40] have shown that it takes 2 h of milling to reach a homogeneous distribution of Pd in MgNi, consistent with the finding of the present study. 3.3. Electrochemical measurements 3.3.1. Hydrogen discharge capacities Fig. 3A shows the cycling discharge capacities of Mg50 Ti50 –3.3 at.% Pd milled for various times. Typically, 3 ± 1 cycles are required before reaching the maximum discharge capacity, independently of the milling duration in presence of Pd. As shown previously, this activation period might be related to an irreversible structural transition from the initial bcc/hcp phase mixture to an fcc phase induced by the hydrogenation process [8,9]. However, an improvement of the charge-transfer reaction on the alloy surface associated with the rupture of the native oxide layer and the increase of its effective surface area during the first few cycles can also explain the activation period. A seen in Fig. 3A, the milling time with Pd has a major influence on the electrode discharge capacities. This is more clearly depicted in Fig. 3B where the maximum discharge capacity (Qmax ) of the materials is plotted vs. the milling time with Pd. For short milling times (≤2 h), Qmax increases rapidly. This variation becomes much less marked at longer milling times (2–20 h). A maximum discharge capacity of 480 mAh g−1 is reached after 20 h of milling with Pd. It can also be noted in Fig. 3A that the electrode capacity decay with cycling is much more rapid (>5% per cycle) for electrodes having the highest Qmax values (>300 mAh g−1 ) whereas electrodes with low Qmax (<150 mAh g−1 ) display almost no capacity loss with cycling. This indicates that the electrode

615

Fig. 3. (A) Cycling discharge capacities of milled Mg50 Ti50 –3.3 at.% Pd electrodes for different milling times with Pd. (B) Evolution of the maximum discharge capacity (Qmax ) of the milled Mg50 Ti50 –3.3 at.% Pd electrodes as a function of the milling time with Pd.

degradation rate strongly depends on the amount of hydrogen absorbed/desorbed from the active material. A similar observation was made with MgNi-based materials [41,42]. It can be explained by the well-known volume expansion/contraction mechanisms occurring upon hydrogen charging/discharging. This process induces an important mechanical stress into the alloy, leading to the disintegration of the particles into smaller fragments. As indicated before, the particle cracking may activate the electrode during the first cycles by breaking the native oxide layer present at the surface of the particles and by increasing the effective surface area of the electrode but on the other hand, the pulverization of particles could also have a negative effect on the electrode cycle life by increasing the corrosion rate of the alloy. It may also cause the loss of electrical connectivity between active material and graphite particles. When hydrogen content in the material is low, the level of mechanical stress into the particles related to their volume expansion/contraction is assumed to be lower than the threshold value where significant particle disintegration occurs. In these conditions, no rapid capacity decay with cycling is observed. 3.3.2. Charge-transfer and hydrogen diffusion kinetics The electrochemical hydrogenation reaction taking place in the metal hydride electrode is mainly governed by the charge-transfer reaction occurring at the electrode/electrolyte interface: M + H2 O + e− ↔ MHads + OH−

(1)

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Fig. 5. Evolution of the exchange current density (I0 ) of the milled Mg50 Ti50 –3.3 at.% Pd electrodes as a function of the Pd surface concentration.

Fig. 4. (A) Linear micropolarization curves of the fully charged milled Mg50 Ti50 –3.3 at.% Pd electrodes for different milling times with Pd. Scan rate: 0.1 mV s−1 . (B) Evolution of the exchange current density (I0 ) of the milled Mg50 Ti50 –3.3 at.% Pd electrodes as a function of the milling time with Pd.

and by the diffusion of hydrogen in the material: MHads ↔ MHabs

(2)

where Hads and Habs represent hydrogen atoms adsorbed at the surface and absorbed into the alloy M, respectively. In order to understand the effect of palladium on the electrochemical hydrogenation of Mg50 Ti50 , the characteristic parameters of reactions (1) and (2) were investigated as a function of the milling time with Pd. The charge-transfer reaction (1) can be characterized by the exchange current density (I0 ). This parameter is extracted from the micropolarization curves presented in Fig. 4A for Mg50 Ti50 – milled with 3.3 at.% Pd for different times. Measurements were performed on electrodes in fully charged state after cycling up to their maximum H-storage capacity. A good linear relationship between the overpotential  and the current I is observed. Exchange current densities are then determined using the linear form of the Butler–Volmer equation valid at low overpotential [43]: I = I0

F RT

(3)

where I0 is the exchange current density, R the universal gas constant, T the temperature and F is the Faraday constant. The calculated values of I0 at their maximum H-storage capacity for the Mg50 Ti50 materials milled from 5 min to 20 h with Pd are depicted in Fig. 4B. I0 is strongly dependent on the milling time with Pd, increasing rapidly from 2.9 mA g−1 for 5 min to 35.3 mA g−1 after 5 h

of milling and then more slowly to 43.3 mA g−1 for 20 h of milling. Similar evolution with respect to the milling time was observed for Pd surface concentration (see Fig. 2C), suggesting that Pd present at the surface of the active material plays a key role in facilitating the charge-transfer reaction. This is clearly illustrated in Fig. 5, which shows an almost linear relationship between the Pd surface concentration and I0 . It is difficult to compare the exchange current densities from different studies because I0 values are not only influenced by the alloy properties but also by the alloy particle size, the electrode state-of-charge, the temperature, the activation of the electrode, etc. Nevertheless, on the basis of the I0 values (after activation and at fully charged state) compiled by Feng et al. [44], the exchange current densities measured on our Mg50 Ti50 materials with low Pd surface concentration are significantly smaller than those measured on LaNi5 -based materials (I0 ∼ 20–300 mA g−1 ) and Zr-based materials (I0 ∼ 20–100 mA g−1 ). They are also smaller than those measured in the same conditions by our group on MgNi-based compounds (I0 = 62–92 mA g−1 ) [45]. This illustrates the poor intrinsic electrocatalytic activity of the Mg50 Ti50 material for the chargetransfer reaction (1). It might be related to the absence of Ni in this material (in contrast to the other compounds cited above) and to the formation of non-conductive compounds (TiO2 , Mg(OH)2 ) on its surface. This is a major disadvantage considering that a slow charge-transfer kinetic induces poor electrode performance, especially under extreme conditions (e.g., low temperature) [46]. Fig. 6 shows the influence of I0 on the Mg50 Ti50 –Pd electrode discharge efficiency (Deff ), defined as: Deff (%) =

Cx × 100 Cx + C5

(4)

where Cx is the maximum discharge capacity obtained at the discharge current density of x mA g−1 and C5 is the residual discharge capacity extracted at a discharge current density of 5 mA g−1 after the electrode is discharged at x mA g−1 . For x = 20 mA g−1 (curve a), the discharge efficiency increases significantly for I0 values smaller than ∼20 mA g−1 (corresponding to [Pd]surface ∼ 2 at.%). For larger I0 values, the discharge current efficiency stays constant. This suggests that, in these conditions, the charge-transfer reaction is the rate-determining step when the Pd surface concentration is lower than ∼2 at.%. For higher Pd concentration, the hydrogen discharge reaction kinetic is under H-diffusion control or under mixed control. On the other hand, for x = 100 mA g−1 (curve b), Deff is low (<30%) and almost independent of I0 . This may indicate that the charge-transfer reaction is the rate-determining step only

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Fig. 6. Evolution of the discharge efficiency (Deff ) of the milled Mg50 Ti50 –3.3 at.% Pd electrodes as a function of the exchange current density (I0 ) for a discharge current density of 20 mA g−1 (curve a) and 100 mA g−1 (curve b).

when the applied discharge current density (ID ) is lower than the exchange current density. For ID > I0 , the hydrogen discharge reaction kinetic is assumed to be under H-diffusion control or under mixed control. Additional works will be needed to fully understand these issues since hydrogen discharge kinetics and its influence on the electrode discharge capacity depend on numerous parameters (state-of-discharge, cycle number, discharge current density, etc.) [47]. Hydrogen diffusion in the active material can be characterized by the hydrogen diffusion coefficient determined from potential step experiments. Measurements were performed on electrodes in fully charged state after cycling up to their maximum H-storage capacity. Fig. 7A shows the resulting current responses expressed in log (i) vs. discharge time for Mg50 Ti50 –3.3 at.% Pd milled for different time. At sufficiently long discharge time (t > ∼20 min), a slow and linear current decay is observed. From this linear part of the curve, it is possible to estimate the average coefficient of hydrogen diffusion, DH , using the following equation [48]: Log i = Log

 6FD

H

da2



(C0 − Cs ) −

2 DH t 2.303a2

(5)

where C0 is the initial hydrogen concentration in the bulk (considered as uniform), CS is the hydrogen concentration at the electrode surface (considered as constant), d is the density of the material, and a is the average radius of the particles (assumed as spherical). The values of DH /a2 were determined for each sample and, as seen in Fig. 7B, it increases from 3.0 to 5.5 × 10−5 s−1 when the milling duration with Pd increases from 15 min to 20 h. However, this apparent evolution with milling time should be interpreted with caution since powder cracking with cycling (affecting a values) may differ for the various samples. Additionally, when the same experiment was repeated several times, a variation of about ±0.5 × 10−5 s−1 in DH /a2 values was observed for the same material. Thus, if any, the effect of milling with Pd on hydrogen diffusion kinetics seems to be limited and much less important than that observed previously on the charge-transfer kinetics. 3.3.3. Electrochemical pressure-composition isotherms The hydrogenation thermodynamic properties of the materials were characterized on the basis of the electrochemical pressurecomposition isotherms (PCT curves) obtained at room temperature (23 ◦ C) from equilibrium potential measurements in the discharge process. Measurements were performed on electrodes in fully charged state after cycling up to their maximum H-storage capac-

Fig. 7. (A) Chronoamperometric curves of the fully charged milled Mg50 Ti50 –3.3 at.% Pd electrodes for different milling times with Pd. Potential step: +50 mV vs. open circuit potential. (B) Evolution of DH /a2 value of the milled Mg50 Ti50 –3.3 at.% Pd materials as a function of the milling time with Pd.

ity. The hydrogen pressure values (PH2 ) were calculated from the measured equilibrium potentials (Eeq ) according to the following equation [49]:

 Eeq

Hg V vs HgO



=

0 EHg/HgO

RT ln − nF



PH2



P0

= −0.926 − 0.0293 log PH2 (atm)

(6)

Fig. 8A displays the PCT curves for Mg50 Ti50 –3.3 at.% Pd for several milling times. They are expressed as a function of the depth of charge of the electrode rather than as a function of the H content in order to facilitate the comparison between the different curves. A well-defined plateau region (related to the ␣-to-␤ phase transition) is observed for all materials with the exception of the sample milled 30 min with Pd. No plateau related to the presence of a Pd phase (expected around −0.9 V vs. Hg/HgO or 0.1 MPa [31]) is discernible, even after short milling time with Pd. This can be explained by the small amount of pure Pd present in the materials after 15 min of milling as shown previously from XRD and XPS analyses. In addition, from 1 to 20 h of milling, the plateau becomes slightly steeper and narrower. The slope formation as the milling time with Pd increases indicates that the distribution of energy levels for hydrogen becomes wider. In other words, the progressive dissolution of Pd atoms in the Mg50 Ti50 alloy with milling must accentuate the chemical disorder in the material, increasing the number of atom coordination types around hydrogen atoms. For

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Fig. 9. Evolution of the maximum discharge capacity (Qmax ) of the 20 h-milled Mg50 Ti50 –x at.% Pd electrodes as a function of their Pd content.

amount and further works will have to be carried out to clarify this issue. 3.4. Influence of the added Pd amount

Fig. 8. (A) Electrochemical hydrogen pressure-composition isotherms of milled Mg50 Ti50 –3.3 at.% Pd materials for different milling times with Pd (in the discharge process at room temperature). (B) Evolution of the half-plateau equilibrium pressure (PH2 eq ) of the milled Mg50 Ti50 –3.3 at.% Pd materials as a function of the milling time with Pd.

instance, 2Mg1Ti1Pd and 1Mg1Pd2Ti sites may be formed in addition to the 2Mg2Ti sites. More importantly, Fig. 8A shows that the hydrogen plateau pressure increases substantially as the milling duration with Pd increases, especially between 15 and 60 min of milling. This is illustrated in Fig. 8B, which displays the evolution of the half-plateau equilibrium pressure PH2 eq of the materials as a function of the milling time with Pd. For 15 min of milling, PH2 eq is very low, around 10−8 MPa. With further milling, this value quickly rises to about 10−5 MPa after 1 h of milling. Between 1 and 20 h of milling with Pd, the variation of PH2 eq is much less marked; it reaches a maximum of 4.6 × 10−5 MPa after 20 h of milling. These values are similar to those found by Kalisvaart and Notten for Pd activated Mg–Ti based alloys prepared by ball milling (plateau around −0.75 V vs. Hg/HgO) [6]. The plateau pressure rise reflects the destabilization of the hydride with increasing milling time with Pd. Geometric considerations (i.e., unit cell volume decrease) can be excluded to explain this effect since the lattice volumes of the bcc and hcp Mg–Ti phases do not change significantly with milling time (see Rietveld analyses). A possible explanation is that the dissolution of Pd changes the electronic structure (i.e., the M–H and M–M bond strengths) in the alloy [50,51], which may be accentuated by the electronegativity difference between Pd on one side and Mg and Ti atoms on the other side (2.2, 1.31 and 1.54 in Pauling scale, respectively). This may decrease the energy of the hydrogen-sites in the Mg–Ti material and has a beneficial effect on the electrochemical hydrogen discharge. However, the major influence of Pd addition on the Mg–Ti hydride stability stays very surprising considering its low

The influence of the amount of Pd added to Mg50 Ti50 was investigated. Pd was introduced at the beginning of the milling with the elemental Mg and Ti powders and the milling time was set to 20 h. The added Pd content was varied from 0 to 6.6 at.% Pd. Dissolution of Pd in Mg50 Ti50 is assumed to be complete for all samples since, as before, the diffraction peaks of pure Pd are no longer observed at the end of the milling period (20 h). The evolution of the maximum discharge capacities of the different samples as function of their Pd contents is depicted in Fig. 9. As expected, the electrode discharge capacity increases with the Pd content but no significant improvement is observed for [Pd] > 3.3 at.%, that corresponds to the optimal Pd concentration. 3.5. Palladium substitution Attempts to substitute Pd by others catalysts were conducted. The tested catalysts were Pt, Ni, Nb and V. Pt and Ni were chosen for their good electrocatalytic activity for the electroreduction of water while Nb and V were chosen for their good H-diffusivity properties. For these experiments, 3.3 at.% of catalyst was added to elemental Mg and Ti powders at the beginning of the milling period and the milling time was set to 20 h. All of these materials exhibit maximum discharge capacities lower than 30 mAh g−1 . It must be noted that all of these materials display a low exchange current density (I0 ≤ 5 mA g−1 ). This confirms the key role played by Pd on the efficiency of the charge-transfer reaction and the direct incidence of this reaction on the electrochemical H-storage capacity of the Mg50 Ti50 alloy. 4. Conclusion The effect of adding Pd on the structure and electrochemical hydrogenation properties of ball-milled Mg50 Ti50 alloy has been studied by varying the milling time in presence of Pd. The main results can be summarized as follows: 1. The crystalline structure of the Mg50 Ti50 alloy (bcc and hcp Mg–Ti phase mixture) does not change significantly with the addition of Pd.

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2. The alloying of Pd with pre-milled Mg50 Ti50 occurs very rapidly (few minutes) and is complete after 5 h of milling. 3. The Pd surface concentration on Mg50 Ti50 increases rapidly with the milling time and reaches the nominal Pd concentration after 5 h of milling. 4. The maximum discharge capacity of Mg50 Ti50 –3.3 at.% Pd increases significantly with the milling time, from 35 mAh g−1 for 5 min to 480 mAh g−1 for 20 h of milling. 5. The exchange current density I0 increases rapidly with the milling time and is directly related to the Pd surface concentration. This demonstrates that Pd present at the surface of the active material plays a key role in facilitating the charge-transfer reaction. 6. The effect of milling with Pd on the hydrogen diffusivity in Mg50 Ti50 is limited and less important than on the chargetransfer kinetics. 7. The electrochemical PCT curves reveal a significant destabilization of the hydride as the milling time with Pd increased. Acknowledgments This work has been financially supported by the Natural Sciences and Engineering Research Council (NSERC) and the “Fonds Québecois de la Recherche sur la Nature et les Technologies” (FQRNT). The authors thank C. Maunders and G. Botton (McMaster University) for the transmission electron microscopy analyses. References [1] R.A.H. Niessen, P.H.L. Notten, Electrochem. Solid-State Lett. 8 (2005) A534. [2] P. Vermeulen, R.A.H. Niessen, P.H.L. Notten, Electrochem. Commun. 8 (2006) 27. [3] P. Vermeulen, R.A.H. Niessen, D.M. Borsa, B. Dam, R. Grissen, P.H.L. Notten, Electrochem. Solid-State Lett. 9 (2005) A520. [4] P. Vermeulen, E.F.M.J. van Thiel, P.H.L. Notten, Chem. Eur. J. 13 (2007) 9892. [5] W.P. Kalisvaart, H.J. Wondergem, F. Bakker, P.H.L. Notten, J. Mater. Res. 22 (2007) 1640. [6] W.P. Kalisvaart, P.H.L. Notten, J. Mater. Res. 23 (2008) 2179. [7] S. Rousselot, M.P. Bichat, D. Guay, L. Roué, J. Power Sources 175 (2008) 621. [8] S. Rousselot, M.P. Bichat, D. Guay, L. Roué, ECS Trans. 16 (2009) 91. [9] S. Rousselot, M.P. Bichat, D. Guay, L. Roué, J. Electrochem. Soc., in press. [10] Y.J. Choi, J. Lu, H.Y. Sohn, Z.Z. Fang, J. Power Sources 180 (2008) 491. [11] K. Asano, H. Enoki, E. Akiba, J. Alloys Compd. 478 (2009) 117. [12] K. Asano, H. Enoki, E. Akiba, J. Alloys Compd. 480 (2009) 558. [13] K. Asano, E. Akiba, J. Alloys Compd. 481 (2009) L8. [14] D.M. Borsa, A. Baldi, M. Pasturel, H. Schreuders, B. Dam, R. Griessen, P. Vermeulen, P.H.L. Notten, Appl. Phys. Lett. 88 (2006) 241910.

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