An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys prepared by mechanical milling

JOURNAL OF RARE EARTHS, Vol. 33, No. 8, Aug. 2015, P. 874

An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys prepared by mechanical milling ZHANG Yanghuan (ᓴ㕞ᤶ)1,2, , ZHANG Pengjun (ᓴ᳟‫)ݯ‬1, YUAN Zeming (㹕⋑ᯢ)2, YANG Tai (ᴼ ⋄)2, QI Yan (⼕ ✅)2, ZHAO Dongliang (䍉ᷟṕ)2 (1. Key Laboratory of Integrated Exploitation of Baiyun Obo Multi-Metal Resources, Inner Mongolia University of Science and Technology, Baotou 014010, China; 2. Department of Functional Material Research, Central Iron and Steel Research Institute, Beijing 100081, China) Received 2 February 2015; revised 17 May 2015

Abstract: The nanocrystalline and amorphous Mg2Ni-type electrode alloys with a composition of Mg20xYxNi10 (x=0, 1, 2, 3 and 4) were fabricated by mechanical milling. Effects of Y content on the structures and electrochemical hydrogen storage performances of the alloys were investigated in detail. The inspections of X-ray diffraction (XRD), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) revealed that the substitution of Y for Mg brought on an obvious change in the phase composition of the alloys. The substitution of Y for Mg resulted in the formation of secondary YMgNi4 phases without altering the major phase Mg2Ni when Y content x1. But with the further increase of Y content, the major phase of the alloys changed into YMgNi4 phase. In addition, such substitution facilitated the glass forming of the Mg2Ni-type alloy. The discharge capacities of the as-milled alloys had the maximum values with Y content varying, but Y content with which the alloy yielded the biggest discharge capacity was changeable with milling time varying. The substitution of Y for Mg had an insignificant effect on the activation ability of the alloys, but it dramatically improved the cycle stability of the as-milled alloys. The effect of Y content on the electrochemical kinetics of the alloys was related to milling time. When milling time was 10 h, the high rate discharge ability (HRD), diffusion coefficient of hydrogen atom (D) and charge transfer rate all had the maximum value with Y content increasing, but they always decreased in the same condition when milling time increased to 70 h. Keywords: Mg2Ni-type alloy; Y substitution for Mg; milling time; electrochemical performance; rare earths

Hydrogen is regarded as one of the most promising alternative fuels offering clean energy which will reduce the intemperate use of non-renewable fossil fuel, partly eliminating the emissions of both greenhouse gas and air pollutants[1]. Currently, one of the main challenges for using hydrogen is the lack of efficient methods of storage due to its very low volumetric and gravimetric density at ambient pressure and temperature[2]. Conventional methods of hydrogen storage include liquid hydrogen[3], compressed gas, metal hydrides[4] and sorption on different porous materials like carbon materials or metal-organic frameworks (MOFs)[5]. As far as the main ways of hydrogen storage, metal hydride systems are deemed to be more accurate, efficient and safe, representing the frontiers of technology[6]. Mg2Ni-type metallic hydrides are looked upon as one of the most promising hydrogen storage materials applied in hydrogen fuel cell vehicle or as negative electrodes in Ni-MH batteries[7] because of their major advantages, such as the higher gaseous hydrogen absorption capacity (3.6 wt.%) for Mg2NiH4 and the larger theoretical electrochemical capacity (about 1000 mAh/g)[8,9]. However, the practical applications of

Mg2Ni-type alloy for Ni-MH batteries are full of challenges because of its sluggish hydriding/dehydriding kinetics[10], low electrochemical discharge capacity relative to the theoretical value at room temperature[11] and poor cycle stability in alkaline solution. Various attempts, particularly mechanical alloying (MA)[12], alloying with other elements[13], adding catalyst[14] and melt spinning[15], have been adopted to overcome these drawbacks and there has been a remarkable improvement in the desorption kinetics[16]. It has come to light that the specific capacity and hydriding/dehydriding kinetics of hydride electrode materials depend on their chemical composition and crystalline structure. Especially, the addition of catalytic elements like transition metals, rare-earth (RE) metals and transition metal oxides decreases the stability of the hydride and facilitates hydrogen desorption[13]. Moreover, nanocrystalline and amorphous Mg-based hydrogen storage alloys show a higher hydrogen absorption capacity at low temperatures and ambient pressure, and better kinetics of hydriding and dehydriding compared with their bulk counterparts[17]. As reported by Kumar et al.[18],

Foundation item: Project supported by the National Natural Science Foundation of China (51161015, 51371094) * Corresponding author: ZHANG Yanghuan (E-mail: [email protected]; Tel.: +86-10-62183115) DOI: 10.1016/S1002-0721(14)60499-3

ZHANG Yanghuan et al., An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys …

ultra-fine microstructure (sub 100 nm range) decreases the absorbing/desorbing hydrogen temperature of Mg2Ni alloy by 100 K, namely from 573 to 473 K. Different methods can be applied to produce such like nanocrystalline Mg-based alloys. Among them, melt spinning is recognized to be very suitable to yield high purity amorphous and/or nanocrystalline alloys with a very homogeneous element distribution at high production rate and low processing costs. Mechanical milling (MM), although having some insurmountable disadvantages in processing by powder metallurgy such as the surface contamination and the time-consuming, is considered to be quite an appropriate technique for producing amorphous and nanocrystalline Mg2Ni alloys with different compositions[19]. In our previous works, a series of Mg2Ni-type alloys were fabricated by melt spinning and their structures and electrochemical performances were investigated[20, 21]. In the present work, the Mg2Ni-type Mg20xYxNi10 (x=04) alloys were prepared by mechanical milling, and the effects of Y content on the structures and electrochemical performances of the alloys were investigated in detail.

1 Experimental The experimental alloys with the chemical composition of the Mg20xYxNi10 (x=04) were prepared using a vacuum induction furnace under a helium atmosphere at a pressure of 0.04 MPa to prevent Mg from volatilizing. The molten alloy was poured into a copper mould, thus a cast ingot was obtained. A part of the as-cast alloys was mechanically crushed into powder with a diameter of about 50 m. Then, the prepared powder was mechanically milled by a planetary-type mill in an argon atmosphere to prevent the powders from being oxidized during ball milling. The samples were handled in a glove box under Ar atmosphere. Cr-Ni stainless steel balls and the powders with a weight ratio of 35:1 were put into Cr-Ni stainless steel vials together. The milling speed was 135 r/min and the duration time was 10 and 70 h, respectively. In order to being convenient for description, the alloys were denoted with Y content as Y0, Y1, Y2, Y3 and Y4, respectively. The phase structures of the as-cast and milled alloys were determined by XRD (D/max/2400). The diffraction, with the experimental parameters of 160 mA, 40 kV and 10 (°)/min respectively, was performed with Cu K1 radiation filtered by graphite. The powder samples of the as-milled alloys were observed by high resolution TEM (JEM-2100F, operated at 200 kV) and their crystalline states were determined by electron diffraction (ED). The mixture of the alloy powder and carbonyl nickel powder in a weight ratio of 1:4 was cold pressed under a pressure of 35 MPa into round electrode pellet with a

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diameter of 15 mm whose total weight is 1 g. The electrochemical performances were measured at 303 K by using a tri-electrode open cell consisting of a working electrode (the metal hydride electrode), a sintered Ni(OH)2/NiOOH counter electrode as well as a Hg/HgO reference electrode, which were immersed in 6 mol/L KOH electrolyte. The voltage between the negative electrode and the reference one was defined as the discharge voltage. In every cycle, the alloy electrode was first charged with a constant current density; after resting for 15 min, it was discharged at the same current density to cut-off voltage of –500 mV. To determine the electrochemical kinetics of the alloy electrodes, the electrochemical impedance spectra (EIS) of the alloys were measured at 303 K by using an electrochemical workstation (PARSTAT 2273). Prior to measurement, several electrochemical charging and discharging cycles were carried out to activate the materials. The fresh electrodes were fully charged and then rested for 2 h up to the stabilization of the open circuit potential. The EIS of the alloy electrodes were measured at 50% depth of discharge (DOD), frequency range from 10 kHz to 5 mHz, amplitude of signal potentiostatic or galvanostatic measurements being 5 mV and the number of points per decade of frequencies being 60. For the potentiostatic discharge, the test electrodes in the fully charged state were discharged at 500 mV potential steps for 5000 s on the electrochemical workstation using electrochemistry corrosion software (CorrWare).

2 Results and discussion 2.1 Microstructure characteristics The phase components and structure characteristics of the as-milled Mg20xYxNi10 (x=04) alloys are subjected to XRD detections, just as depicted in Fig. 1, from which it is found that the substitution of Y for Mg brings on an obvious change in the phase composition of the as-milled alloys. Y0 alloys milled for 10 and 70 h are Mg2Ni (JCPDS 04-004-6583) single phase. When Y content x=1, the substitution of Y for Mg gives rise to the formation of the secondary phases YMgNi4 (JCPDS 01-072-9165) without altering major phase Mg2Ni. But with Y content further increasing, the major phase of the as-milled alloys changes from Mg2Ni into YMgNi4. The lattice parameters and unit cell volumes of the major phases in the alloys are listed in Table 1. It can be seen that a, c, V of the Mg2Ni phase and a, V of YMgNi4 phases in the alloys increase with increasing Y content x, which is mainly ascribed to the fact that the atom radius of Y is larger than that of Mg. Moreover, we also find that with milling time prolonging from 10 to 70 h, the parameters of major phases and the widths of diffraction peaks are also increased. It is mainly because of the grain refinement and

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Fig. 1 XRD profiles of the as-milled Mg20xYxNi10 (x=04) alloys (a) Milling for 10 h; (b) Milling for 70 h Table1 Lattice parameters and unit cell volumes of the major phase of the as-milled Mg20xYxNi10 (x=04) alloys 10 h Alloys Major phase

a/nm

2.2 Electrochemical performances

70 h c/nm V/nm3

Major phase

a/nm

observing Debye-Scherrer rings.

c/nm V/nm3

Y0

Mg2Ni 0.52071 1.3223 0.31050

Mg2Ni 0.52235 1.3279 0.31378

Y1

Mg2Ni 0.52334 1.3280 0.31499

Mg2Ni 0.52369 1.3273 0.31523

Y2

YMgNi4 0.70926

–

0.35681 YMgNi4 0.71078

–

0.35909

Y3

YMgNi4 0.71046

–

0.35860 YMgNi4 0.71212

–

0.36113

Y4

YMgNi4 0.71050

–

0.35866 YMgNi4 0.71261

–

0.36188

the creation of crystal defects during ball milling. The morphology characteristics of the as-milled Mg20xYxNi10 (x=04) alloys were examined by TEM, just as illustrated in Fig. 2. It is obvious that when milling time is 10 h, the Y0 and Y1 alloys have a nearly entire nanocrystalline structure, while Y3 alloy exhibits a visible nanocrystalline and amorphous structures, indicating that the substitution of Y for Mg facilitates the glass forming of the Mg2Ni alloy during milling, which is supported by observing Debye-Scherrer rings. A similar conclusion was reported by Teresiak et al.[22]. Meanwhile, we find that when milling time is prolonged from 10 to 70 h, the grain sizes of the alloys markedly decreases and the disordered degree of their microstructures obviously increases, meaning that a large amount of strain energy is stored, which leads to non-stabilization of the lattice, yielding the fine grain sizes. Moreover, after careful checking, we can clearly find an amorphous phase in the as-milled (70 h) Y1 and Y3 alloys, which is confirmed by

2.2.1 Activation capability, discharge potential and discharge capacity The activation capability was evaluated by the number of charging-discharging cycles required for attaining the greatest discharge capacity through repeated chargedischarge process at a constant current density. The fewer the number of charging-discharging cycle is, the better the activation performance will be. Easy to be activated is necessary for the alloy electrode applied in Ni-MH battery. The variations of the discharge capacities of the as-milled Mg20xYxNi10 (x=04) alloys with cycle number are presented in Fig. 3. Apparently, all the alloys attain their maximum discharge capacities at the first charging-discharging cycle, exhibiting the excellent activation capability. The variation of Y content does not affect the activation capability of the alloys. Shown in Fig. 4 are the discharge potential curves of the as-milled Mg20xYxNi10 (x=04) alloy electrodes with a current density of 40 mA/g at the first charging/discharging cycle. The discharge potential characteristic, which directly determines the stability of the output power, is a very important performance of the alloy electrode and is characterized by the potential plateau of the discharge curve of the alloy. The longer and more horizontal the discharge potential plateau is, the better the discharge potential characteristics of the alloy will be. We note that when milling time is 10 h, the discharge potential char-

ZHANG Yanghuan et al., An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys …

Fig. 2 TEM micrographs and ED patterns of the as-milled Mg20xYxNi10 (x=04) alloys (a), (b) and (c) Y0, Y1 and Y3 alloys milled for 10 h; (d), (e) and (f) Y0, Y1 and Y3 alloys milled for 70 h

Fig. 3 Evolution of the discharge capacity of the as-milled Mg20xYxNi10 (x=04) alloys with cycle number (a) Milling for 10 h; (b) Milling for 70 h

Fig. 4 Discharge potential curves of the as-cast and milled Mg20xYxNi10 (x=04) alloys (a) Milling for 10 h; (b) Milling for 70 h

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acteristic of the alloys is always ameliorated with Y content increasing, but when milling time is 70 h, the discharge potential characteristic of the alloys is improved till Y content x=1 and then it markedly worsens with Y content growing. Lai et al.[23] considered that the discharge potential of an alloy electrode be closely related to internal resistance of the battery, including ohmic internal resistance and polarization resistance, which mainly depends on the diffusion of hydrogen atoms in the alloy and it reduces with the increase of the diffusion coefficient of the hydrogen atom. The influence of Y content and milling time on the diffusion coefficient of hydrogen atom will be discussed in the next section. Furthermore, it is found that the substitution of Y for Mg and mechanical milling engender an obvious effect on the discharge capacity of the alloys. Based on the data in Fig. 4, the relationship between the discharge capacity and the Y content can be established easily, as described in Fig. 5. It is evident that the discharge capacities of the alloys have the maximum values with Y content varying, but Y content with which the alloys obtain the maximum discharge capacity is different with milling time altering. To be specific, Y content corresponding to the maximum discharge capacity is x=3 as milling time is 10 h, whereas it is x=1 when milling time is 70 h. The discharge capacities of the alloys have maximum values with Y content varying, meaning that Y substitution for Mg gives rise to the beneficial and detrimental impacts on the discharge capacity of the alloy. The positive action is ascribed to the decreased thermal stability of the Mg-based hydride caused by Y substitution for Mg[24], enhancing the electrochemical discharge ability of the Mg-based alloy. And the negative function is most likely associated with the facilitated glass forming by Y substitution for Mg due to the fact that an amorphous phase has a lower discharge capacity than a crystal counterpart. In addition, it is noteworthy that when the Y content x2, the substitution of Y for Mg leads to the major phase changing from Mg2Ni to YMgNi4, which necessarily makes the discharge capacity decrease.

Fig. 5 Evolution of the discharge capacity of the as-milled Mg20xYxNi10 (x=04) alloys with Y content

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2.2.2 Electrochemical cycle stability Whether or not a kind of alloy can be applied as a negative electrode material of Ni-MH battery depends on its electrochemical cyclic stability, which is characterized by capacity retaining rate (Sn), being defined as Sn=Cn/ Cmax×100%, where Cmax is the maximum discharge capacity and Cn is the discharge capacity of the nth chargedischarge cycle at a current density of 40 mA/g, respectively. If an alloy electrode can maintain a higher capacity retention rate after being repeated charge-discharge cycles, it means that the alloy electrode owns superior electrochemical cyclic stability and is more appropriate to be the negative materials for Ni-MH batteries. The variations of the Sn values of the as-milled Mg20xYxNi10 (x=04) alloys with cycle number are described in Fig. 6. It is evident that the Sn values of the alloys decline with cycle number increasing. We note that the degradation rate of discharge capacity of the as-milled alloys dramatically decreases with Y content increasing, suggesting that the substitution of Y for Mg makes a positive contribution to the cycle stability of the alloys. More specifically, the S20 value of the alloy is enhanced from 38.3% to 80.2% for the alloy milled for 10 h and from 33.9% to 74.7% for the alloy milled for 70 h by increasing Y content from 0 to 4. Furthermore, a careful comparison finds that, for a fixed Y content, increasing milling time from 10 to 70 h makes the capacity retaining rate of the alloy clearly decrease, meaning that prolong-

Fig. 6 Evolution of the capacity retaining rates (Sn) of the asmilled Mg20xYxNi10 (x=04) with cycle number (a) Milling for 10 h; (b) Milling for 70 h

ZHANG Yanghuan et al., An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys …

ing milling time impairs the cycle stability of the alloys. Here, some elucidations can be given about the effects of Y content and milling time on the cycle stability of the alloys. As is well known, the major cause of leading to the sharp degradation of the discharge capacity of Mg2Ni alloy is forming and thickening of Mg(OH)2 layer on the surface of the alloy electrode, which hinders hydrogen atoms from diffusing in or out, in alkaline solution[25]. Moreover, hydrogen storage material suffers from an inevitable volume change during the charge/discharge process resulting in the cracking and pulverizing of the alloy which, in turn, makes the surface of the material apt to be oxidized. Moreover, the capacity degradation of the alloy is convinced to be strongly related to the formation of lattice strain containing lattice defects during the milling[26]. Our experimental results also confirm that the major cause of leading to the capacity degradation of the alloy is from corrosion, as depicted in Fig. 7 from which it is found that the size of the alloy particles has no obvious change after electrochemical cycle, suggesting that the pulverization of alloy particles scarcely takes place in the process of the electrochemical cycle. However, a rough and flocculent layer can clearly be seen on the surface of the alloy particles after electrochemical cycling, which is determined to be magnesium hydroxide by XRD detection, as presented in Fig. 7(c). Apparently, the formation of the rough and flocculent layer is an important reason of leading to the capacity deterioration of the alloy. The beneficial effect of Y substitution for Mg on the cycle stability is ascribed to the following several aspects. Firstly, the partial substitution of Y for Mg can enhance the resistance of the alloy against corrosion in alkaline solution[27], which is considered to be associated with forming a dense and protective yttrium oxide on the alloy surface that suppresses further oxidation of Mg and induces a Ni enriched layer on the alloy surface[28], thereby increasing the cycle life of the alloy. Secondly, the rare earth elements (La, Nd, Sm and Y) can be dissolved in the Mg2Ni alloy, which enlarges the cell volume of the Mg2Ni alloy easily[29] and decreases the ratios

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of expansion/contraction in the process of hydrogen absorption/desorption, thus enhancing the anti-pulverization capability. In respect of the adverse impact caused by mechanical milling, it is convinced to be caused by forming the nanocrystalline structure and producing the lattice strain resulted from milling because intergranular corrosion is unavoidable[30]. 2.2.3 Electrochemical kinetics Keeping a high discharge capacity even during the process of charge-discharge cycles with a big current density is necessary for the practical application of alloy electrode in Ni-MH battery, especially for power battery. It is well known that a decrease in the discharge capacity of the alloy electrode with current density increasing is unavoidable. Usually, the electrochemical kinetics of an alloy electrode is symbolized by its HRD which is defined as: HRD=Ci/C40×100%, where Ci and C40 are the maximum discharge capacities of the alloy electrode charged-discharged at the current densities of i and 40 mA/g, respectively. The evolutions of HRDs of the asmilled Mg20xYxNi10 (x=04) alloys with the discharge current density are presented in Fig. 8, from which it is found that the HRD values of the electrode alloys reduce with the current density rising, and the substitution of Y for Mg and milling time have an evident effect on the HRD of the alloys. On the basis of the data in Fig. 8 at a current density of 200 mA/g (i=200 mA/g), the relationships between the HRD of the as-milled alloys and Y content can be established, as inserted in Figs. 8(a) and (b), respectively. We note that the variation trends of the HRD of the alloys with Y content are different when milling time is 10 and 70 h. For milling for 10 h, the HRD of the alloy has a maximum value with Y content increasing, but for milling for 70 h, it always declines in the same condition. The electrochemical hydriding/dehydriding reaction taking place at the hydrogen storage electrode in an alkaline solution during charging and discharging can be described as follows:

Fig. 7 SEM morphologies together with typical XRD patterns of the as-milled (70 h) Y1 alloys before and after electrochemical cycle (a) Before cycling; (b) After cycling; (c) XRD pattern after cycling

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Fig. 8 Evolution of the HRDs of the as-cast and milled Mg20xYxNi10 (x=04) alloys with current density (a) Milling for 10 h; (b) Milling for 70 h

M

x

H2O 

x



e l MH x 

x

OH



(1)

2 2 2 where M is the hydrogen storage alloy. Eq. (1) indicates that, when the alloy electrode is charged in KOH solution, hydrogen atoms on the alloy-electrolyte interface diffuse into bulk alloy and then store themselves in the metallic lattice in the form of hydride. In the process of discharging, the hydrogen stored in the bulk alloy diffuses toward the alloy electrode surface where it is oxidized. It indicates that electrochemical hydrogen storage kinetics of the alloy electrode is dependent on the hydrogen diffusion capability in the alloy bulk and the charge-transfer rate on the surface of an alloy electrode. Hence, it is very necessary to investigate the effects of Y content and milling time on the charge-transfer rate and the diffusion ability of hydrogen atoms to reveal the mechanism of the electrochemical kinetics of the alloy impacted by Y content and milling time. Hydrogen diffusion ability is characterized by the hydrogen diffusion coefficient, which can be derived by means of the semilogarithmic curves of anodic current versus working duration of an alloy electrode, as demonstrated in Fig. 9. Based on the White’s model[31], the diffusion coefficient of the hydrogen atoms in the bulk of the alloy could be calculated easily through the slope of the linear region of the corresponding plots according to the following formulae:

Fig. 9 Semilogarithmic curves of anodic current vs. time responses of the as-milled Mg20xYxNi10 (x=04) alloys with current density (a) Milling for 10 h; (b) Milling for 70 h

lg i

2 § 6 FD (C  C ) ·   D t 0 s ¸ © da 2 ¹ 2.303 a 2

lg ¨ r

(2)

2

D



2.303a d lg i 

2

(3)

dt

where i is the diffusion current density (A/g), D is the hydrogen diffusion coefficient (cm2/s), C0 is the initial hydrogen concentration in the bulk of the alloy (mol/cm3), Cs is the hydrogen concentration on the surface of the alloy particles (mol/cm3), a is the alloy particle radius (cm), d is the density of the hydrogen storage alloy (g/cm3), t is the discharge time (s), respectively. The evolutions of the D values of the alloys derived by Eq. (3) with Y content are also presented in Figs. 9(a) and (b), respectively. Clearly, the D values of the alloys first augment and then decline with Y content increasing for milling 10 h, while it always decreases in the same condition for milling 70 h, which is very similar to the variation tendencies of the HRDs of the alloys with Y content, indicating that the diffusion ability of hydrogen atoms is a very important factor affecting the electrochemical kinetics of the alloy. Northwood et al.[32] considered that the diffusion coefficient of hydrogen atoms in the metallic lattices is dominated by the strength of the metal-hydrogen interaction and the structure of the alloy[32]. The positive contribution of Y substitution for Mg to hydrogen diffusion is ascribed to following two as-

ZHANG Yanghuan et al., An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys …

pects. Kalinichenka et al.[33] confirmed that the partial substitution of rare earth elements (La, Ce, Pr, Nd, Y and Sm) for Mg and (Cu, Fe, Al, Cr, Co, Mn) for Ni in Mg2Ni alloy can weaken the bond between Mg and H atoms, consequently bringing on an increase in hydrogen diffusion. Also, Cui et al.[34] reported that an increase in the lattice constants and cell volume facilitates to decrease the diffusion activation energy of hydrogen atoms, prompting hydrogen diffusion. Evidently, it can also be found from Fig. 9 that the substitution of Y for Mg causes a negative impact on the hydrogen diffusion, which is ascribed to the facilitated glass forming by Y substitution for Mg since an amorphous phase impairs the diffusion ability of hydrogen atoms[10]. It is favorable and unfavorable impacts caused by Y substitution for Mg that give rise to the diffusion coefficient of hydrogen atoms in the as-milled (10 h) alloys having a maximum value with Y content varying (Fig. 9(a)). Differing milling for 10 h, the milling for 70 h makes the much more amorphous phase form in the alloy due to an indisputable fact that the amount of amorphous phase is proportional to milling time. Therefore, it is easy to understand that the diffusion coefficient of hydrogen atoms of the as-milled (70 h) alloys always decreases with Y content increasing. As for charge-transfer ability of the surface of an alloy electrode, it can be qualitatively evaluated through measuring EIS by means of the Kuriyama’s model[35]. Illustrated in Fig. 10 are the EIS curves of the as-milled Mg20xYxNi10 (x=04) alloys from which it is found that each EIS has two distorted capacitive loops at the high

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and middle frequency regions separately as well as a line at the low frequency region, which very well expresses the electrochemical process of the alloy electrode. Among them, the smaller semicircle in the high frequency region is regarded to reflect the contact resistance between the alloy powder and the conductive material, and the larger one in the middle frequency region corresponds to the charge-transfer resistance on the alloy surface while the straight line in the low frequency represents the atomic hydrogen diffusion in the alloy. In view of this, the charge-transfer ability can be evaluated easily, namely the larger the radius of the semicircle in the middle frequency region is, the higher the charge-transfer resistance of the alloy electrode will be. We note that the variation trend of the radii of the large semicircles of the alloys in the middle frequency region with Y content varying is dependent on the milling time. When milling time is 10 h, the radii of the large semicircles of the alloys in the middle frequency region first shrinks and then expands with Y content increasing, but it obviously enlarges for milling for 70 h in the same condition. Kleperis et al.[36] considered that the charge-transfer step is determined by both crystallographic and electronic structure jointly. The variation of the alloy compositions on the alloy surface results in an evident impact on the valence electron configurations, which essentially controls the charge-transfer reaction, namely the hydrogen dissociative reaction[37]. The positive contribution of Y substitution for Mg to the charge transfer is considered to be associated with the improved corrosion resistance by substituting Mg with Y which suppresses further oxidation of Mg and induces a Ni enriched layer on the alloy surface[28], enhancing the electrocatalytic activity of the surface of the alloy electrode. The adverse action of Y substitution for Mg on the charge transfer is ascribed to facilitated glass forming by Y substitution for Mg because an amorphous phase can strongly prohibit the pulverization of the alloy during charge-discharge cycle[38], reducing available new surface of the alloy electrode and lessening charge transfer rate at the alloy-electrolyte interface. With Y content increasing (from 2 to 4) and milling time prolonging (from 10 h to 70 h), the glass forming ability of the alloys is dramatically enhanced. So far, it seems to be needless to say that the charge transfer rate of the surface of the alloy electrode decreases with Y content increasing.

3 Conclusions

Fig. 10 EIS of the as-milled Mg20xYxNi10 (x=04) alloys (a) Milling for 10 h; (b) Milling for 70 h

(1) The Mg2Ni-type Mg20xYxNi10 (x=04) alloys with nanocrystalline and amorphous structures were fabricated by mechanical milling, and the effects of Y content on the structures and electrochemical performances of the alloys were investigated. The results discovered that the substitution of Y for Mg brought on the formation of

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the secondary phases YMgNi4. And as Y content x2, the substitution of Y for Mg gave rise to the major phase of the alloy changing from Mg2Ni into YMgNi4. Furthermore, such substitution facilitated the glass forming in the as-milled alloy. (2) The electrochemical measurement indicated that all the experimental alloys exhibited excellent activation capability. The discharge capacities of the alloys had the maximum values with Y content varying, but Y content with which the alloys obtained the maximum discharge capacity was variable with milling time changing. (3) Furthermore, the discharge abilities of the as-milled alloys were determined by hydrogen diffusion capability in the alloy bulk and the charge-transfer rate on the surface of an alloy electrode. The variations of electrochemical kinetics of the alloys with Y content were dependent on milling time. When milling time was 10 h, the electrochemical kinetics, including HRD, diffusion coefficient (D) and charge transfer rate, first increased and then decreased with Y content growing, but when milling time was 70 h, the electrochemical kinetics always decreased in the same condition.

References: [1] Zheng J Y, Liu X X, Xu P, Liu P F, Zhao Y Z, Yang J. Development of high pressure gaseous hydrogen storage technologies. Int. J. Hydrogen Energy, 2012, 37(1): 1048. [2] Kowalczyk P, Holyst R, Terrones M, Terrones H. Hydrogen storage in nanoporous carbon materials: myth and facts. Phys. Chem. Chem. Phys., 2007, 9: 1786. [3] Aceves S M, Berry G D, Martinez-Frias J, Espinosa-Loza F. Vehicular storage of hydrogen in insulated pressure vessels. Int. J. Hydrogen Energy, 2006, 31(15): 2274. [4] De Lima G F, Garroni S, Baró M D, Suriach S, Kiminami C S, Botta W J, Jorge Jr A M. Hydrogen storage properties of 2Mg-Fe mixtures processed by hot extrusion: Influence of the extrusion ratio. Int. J. Hydrogen Energy, 2012, 37(20): 15196. [5] Hirscher M, Becher M, Haluska M, von Zeppelin F, Chen X H, Dettlaff-Weglikovska U, Roth S. Are carbon nanostructures an efficient hydrogen storage medium? J. Alloys Compd., 2003, 356-357: 433. [6] Sakintuna B, Lamati-Darkrim F, Hirscher M. Metal hydride materials for solid hydrogen storage: A review. Int. J. Hydrogen Energy, 2007, 32(9): 1121. [7] Hsu F, Lin C, Lee S, Lin C, Bor H. Effect of Mg3MnNi2 on the electrochemical characteristics of Mg2Ni electrode alloy. J. Power Sources, 2010, 195(1): 374. [8] Huang L W, Elkedim O, Nowak M, Jurczyk M, Chassagnon R, Meng D W. Synergistic effects of multiwalled carbon nanotubes and Al on the electrochemical hydrogen storage properties of Mg2Ni-type alloy prepared by mechanical alloying. Int. J. Hydrogen Energy, 2012, 37(2): 1538. [9] Schlapbach L, Züttel A. Hydrogen-storage materials for mobile applications. Nature, 2001, 414(6861): 353.

JOURNAL OF RARE EARTHS, Vol. 33, No. 8, Aug. 2015 [10] Xie D H, Li P, Zeng C X, Sun J W, Qu X H. Effect of substitution of Nd for Mg on the hydrogen storage properties of Mg2Ni alloy. J. Alloys Compd., 2009, 478(1-2): 96. [11] Anik M, Gasan H, Topcu S, Akay I, Aydinbeyli N. Electrochemical hydrogen storage characteristics of Mg1.5Al0.5xZrxNi (x=0, 0.1, 0.2, 0.3, 0.4, 0.5) alloys synthesized by mechanical alloying. Int. J. Hydrogen Energy, 2009, 34(6): 2692. [12] Song M Y, Kwak Y J, Shin H S, Lee S H, Kim B G. Improvement of hydrogen-storage properties of MgH2 by Ni, LiBH4, and Ti addition. Int. J. Hydrogen Energy, 2013, 38(4): 1910. [13] Li P, Zhang J, Zhai F Q, Ma G, Xu L, Qu X H. Effect of annealing treatment on the anti-pulverization and anticorrosion properties of La0.67Mg0.33Ni2.5Co0.5 hydrogen storage alloy. J. Rare Earths, 2015, 33(4): 417. [14] Wu Z, Chen L X, Xiao X Z, Fan X L, Li S Q, Wang Q D. Influence of lanthanon hydride catalysts on hydrogen storage properties of sodium alanates. J. Rare Earths, 2013, 31(5): 502. [15] Zhang Q A, Jiang C J, Liu D D. Comparative investigations on the hydrogenation characteristics and hydrogen storage kinetics of melt-spun Mg10NiR (R=La, Nd and Sm) alloys. Int. J. Hydrogen Energy, 2012, 37(14): 10709. [16] Ohara R, Lan C H, Hwang C S. Electrochemical and structural characterization of electroless nickel coating on Mg2Ni hydrogen storage alloy. J. Alloys Compd., 2013, 580: S368. [17] Wang Y, Qiao S Z, Wang X. Electrochemical hydrogen storage properties of ball-milled NdMg12 alloy with Ni powders. Int. J. Hydrogen Energy, 2008, 33(3): 1023. [18] Hima Kumar L, Viswanathan B, Srinivasa Murthy S. Hydrogen absorption by Mg2Ni prepared by polyol reduction. J. Alloys Compd., 2008, 461(12): 72. [19] Ebrahimi-Purkani A, Kashani-Bozorg S F. Nanocrystalline Mg2Ni-based powders produced by high-energy ball milling and subsequent annealing. J. Alloys Compd., 2008, 456(1-2): 211. [20] Zhang Y H, Qi Y, Zhao D L, Guo S H, Ma Z H, Wang X L. An investigation of hydrogen storage kinetics of meltspun nanocrystalline and amorphous Mg2Ni-type alloys. J. Rare Earths, 2011, 29(1): 87. [21] Zhang Y H, Wang H T, Yang T, Zhai T T, Zhang G F, Zhao D L. Electrochemical hydrogen storage performances of the nanocrystalline and amorphous (Mg24Ni10Cu2)100–x Ndx (x=0–20) alloys applied to Ni-MH battery. J. Rare Earths, 2013, 31(12): 1175. [22] Teresiak A, Gebert A, Savyak M, Uhlemann M, Mickel C, Mattern N. In situ high temperature XRD studies of the thermal behaviour of the rapidly quenched Mg77Ni18Y5 alloy under hydrogen. J. Alloys Compd., 2005, 398(1-2): 156. [23] Lai W H, Yu C Z. Research survery of improving discharge voltage platform for Ni-MH battery. Battery Bimonthly, 1996, 26(4): 189. [24] Lass E A. Hydrogen storage measurements in novel Mgbased nanostructured alloys produced via rapid solidification and devitrification. Int. J. Hydrogen Energy, 2011, 36(17): 10787.

ZHANG Yanghuan et al., An investigation on electrochemical hydrogen storage performances of Mg-Y-Ni alloys … [25] Meli F, Züttel A, Schtapbach L. Electrochemical and surface properties of low cost, cobalt-free LaNi5-type hydrogen storage alloys. J. Alloys Compd., 1993, 202(1-2): 81. [26] Zaluski L, Zaluska A, Ström-Olesen J O. Nanocrystalline metal hydrides. J. Alloys Compd., 1997, 253-254: 70. [27] Lenain C, Aymard L, Dupont L, Tarascon J-M. A new Mg0.9Y0.1Ni hydride forming composition obtained by mechanical grinding. J. Alloys Compd., 1999, 292(12): 84. [28] Ruggeri S, Roué L, Huot J, Schulz R, Aymard L, Tarascon J M, Properties of mechanically alloyed Mg-Ni-Ti ternary hydrogen storage alloys for Ni-MH batteries. J. Power Sources, 2002, 112(2): 547. [29] Zhang Y H, Li B W, Ren H P, Guo S H, Wu Z W, Wang X L. An investigation on the hydrogen storage characteristics of the melt-spun nanocrystalline and amorphous Mg20xLaxNi10 (x=0, 2) hydrogen storage alloys. Mater. Chem. Phys., 2009, 115(1): 328. [30] Simi i M V, Zduji M, Dimitrijevi R, Nikoli -Bujanovi

Lj, Popovi N H. Hydrogen absorption and electrochemical properties of Mg2Ni-type alloys synthesized by mechanical alloying. J. Power Sources, 2006, 158(1): 730. [31] Zheng G, Popov B N, White R E. Electrochemical determination of the diffusion coefficient of hydrogen through an LaNi4.25Al0.75 eletrode in alkaline aqueous solution. J. Electrochem. Soc., 1995, 142(8): 2695.

883

[32] Feng F, Northwood D O. Hydrogen diffusion in the anode of Ni/MH secondary batteries. J. Power Sources, 2004, 136(2): 346. [33] Kalinichenka S, Röntzsch L, Riedl T, Weißgärber T, Kieback B. Hydrogen storage properties and microstructure of melt-spun Mg90Ni8RE2 (RE=Y, Nd, Gd). Int. J. Hydrogen Energy, 2011, 36(17): 10808. [34] Cui N, Luo J L. Electrochemical study of hydrogen diffusion behavior in Mg2Ni-type hydrogen storage alloy electrodes. Int. J. Hydrogen Energy, 1999, 24(1): 37. [35] Kuriyama N, Sakai T, Miyamura H, Uehara I, Ishikawa H, Iwasaki T. Electrochemical impedance and deterioration behavior of metal hydride electrodes. J. Alloys Compd., 1993, 202(1-2): 183. [36] Kleperis J, Wójcik G, Czerwinski A, Skowronski J, Kopczyk M, Beltowska-Brzezinska M. Electrochemical behavior of metal hydrides. J. Solid State Electrochem., 2001, 5(4): 229. [37] Nobuhara K, Kasai H, Dino W A, Nakanishi H. H2 dissociative adsorption on Mg, Ti, Ni, Pd and La surfaces. Surf. Sci., 2004, 566-568: 703. [38] Zhang Y H, Li B W, Ren H P, Cai Y, Dong X P, Wang X L. Cycle stabilities of the La0.7Mg0.3Ni2.55xCo0.45Mx (M= Fe, Mn, Al; x=0, 0.1) electrode alloys prepared by casting and rapid quenching. J. Alloys Compd., 2008, 458(1-2): 340.