Electrochemical characteristics of La0.59Nd0.14Mg0.27Ni3.3 alloy with rhombohedral-type and hexagonal-type A2B7 phases

Accepted Manuscript Electrochemical characteristics of La0.59Nd0.14Mg0.27Ni3.3 alloy with rhombohedral-type and hexagonal-type A2B7 phases Cong Wu, Lu Zhang, Jingjing Liu, Yuan Li, Shuqin Yang, Baozhong Liu, Shumin Han PII:

S0925-8388(16)32960-7

DOI:

10.1016/j.jallcom.2016.09.206

Reference:

JALCOM 39040

To appear in:

Journal of Alloys and Compounds

Received Date: 20 April 2016 Revised Date:

18 September 2016

Accepted Date: 20 September 2016

Please cite this article as: C. Wu, L. Zhang, J. Liu, Y. Li, S. Yang, B. Liu, S. Han, Electrochemical characteristics of La0.59Nd0.14Mg0.27Ni3.3 alloy with rhombohedral-type and hexagonal-type A2B7 phases, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.09.206. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Electrochemical characteristics of La0.59Nd0.14Mg0.27Ni3.3 alloy with rhombohedral-type and hexagonal-type A2B7 phases Cong Wua,b, Lu Zhanga,b, Jingjing Liua,b, Yuan Lia,b, Shuqin Yanga,b, Baozhong Liuc, Shumin

a

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Hana,b∗ State Key Laboratory of Metastable Materials Science and Technology, Yanshan University,

Qinhuangdao 066004, PR China

College of Environmental and Chemical Engineering, Yanshan University, Qinhuangdao 066004,

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b

c

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PR China

School of Materials Science & Engineering, Henan Polytechnic University, Jiaozuo 454000, PR

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China

*

Corresponding author: Tel.: +86-335-8074648, Fax: +86-335-8074648.

E-mail address: [email protected]. 1

ACCEPTED MANUSCRIPT Abstract A La0.59Nd0.14Mg0.27Ni3.3 alloy with 72 wt.% Gd2Co7-type (3R-A2B7) and 28 wt.% Ce2Ni7-type (2H-A2B7) phases is prepared by induction melting followed by the

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annealing at 1148 K. Upon further increasing the annealing temperature to 1248 K, the 3R-A2B7 phase transforms into the 2H-A2B7 phase, forming an alloy with a 97 wt.% 2H-A2B7 phase and a 3 wt.% 3R-A2B7 phase. Electrochemical measurement

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shows that the discharge capacity retention of the alloy electrode at the 100th cycle

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increases from 88.5% to 92.4% as the 2H-A2B7 phase increases from 28 wt.% to 97 wt.%. Compared to the 3R-A2B7 phase, the 2H-A2B7 phase shows a stronger ability to resist amorphousiation during dehydrogenation and has better structural stability. Therefore, the transformation from the 3R-A2B7 phase to the 2H-A2B7 phase

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stabilizes the structure of the alloy against amorphization and oxidation, thus improving the cyclic stability of the alloy. Electrochemical pressure–composion (P–C ) isotherms contain two discharge plateaus; the higher plateau corresponds to the

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3R-A2B7 phase and the lower one is associated with the 2H-A2B7 phase. As the

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2H-A2B7 phase increases, the high-rate dischargeability of the alloy electrodes increases from 56.8% to 66.3%.

Keywords: Nickel/metal hydride battery; Re–Mg–Ni-based hydrogen storage alloy; Super-stacking structure; Electrochemical properties; Cyclic stability

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ACCEPTED MANUSCRIPT 1. Introduction Rare earth (Re)–Mg–Ni-based hydrogen storage alloys have been the focus of research interest for applications as negative electrode materials for nickel/metal

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hydride (Ni/MH) batteries because of their high energy density compared with conventional AB5-type alloys and their fast activation properties compared with conventional AB2-type alloys [1–3]. These Re–Mg–Ni-based hydrogen storage alloys

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are composed of layered structures in the form of AB3, A2B7, A5B19 and AB4 (A = Re,

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Mg and B = Ni) with general formulas of m[ReMgNi4]·n[ReNi5] (where m and n are positive integers); in these structures, [ReNi5] (CaCu5-type) and [ReMgNi4] (Laves-type) subunits stack along the c axis alternatively according to certain combinations [3]. Among them, the A2B7-type Re–Mg–Ni-based alloy has been

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studied widely due to its superior overall electrochemical properties [4,5]. Cao et al. reported that La0.95Sm0.66Mg0.40Ni6.25Al0.42Co0.32 A2B7-type alloy with 96.3 wt.% La2Ni7 phase and 3.7 wt.% LaNi5 phase showed good electrochemical characteristics,

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with a maximum discharge capacity of 372 mAh g–1 and 80% discharge capacity

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retention after 239 cycles [6]. With respect to stacking type, the A2B7 phase has a crystal isomer with both Ce2Ni7-type hexagonal (2H) structure and Gd2Co7-type rhombohedral (3R) structure. In the 2H- and 3R-type structures, each unit cell consists of two and three blocks, which corresponds to 2:1 stacking of [ReNi5] layers and a [ReMgNi4] layer along the c axis in each block, respectively. Moreover, the Laves subunits and the atomic numbers in the 2H and 3R structures layers are different [7,8]. The structural differences between the 2H-type and 3R-type phases result in the

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ACCEPTED MANUSCRIPT different electrochemical characteristics of these two type phases. Wang et al. found that the increasing abundance of Gd2Co7-type phase abundance led to the remarkable increase in discharge capacity of the La0.72Nd0.08Mg0.2Ni3.4Al0.1 alloy electrode [9].

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Zhang et al. reported that the cyclic stability of the La1.5Mg0.5Ni7.0 alloy enhanced with the Ce2Ni7-type phase appearing [10]. Due to the few differences in phase composition between 2H-A2B7 phase and 3R-A2B7 phase, the obtained A2B7-type

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Re–Mg–Ni-based alloys are usually composed of 3R-A2B7 and 2H-A2B7 phases.

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Recently, with the deepening of theoretical study and development of preparation technology, La–Mg–Ni-based alloys with single 2H-A2B7 phase have been obtained [11,12]. Denys et al. prepared La1.5Mg0.5Ni7 alloy with 2H-type (La,Mg)2Ni7 single phase using a stepwise sintering method [11]. Our group prepared the La1.6Mg0.4Ni7

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alloy with single 2H-A2B7 phase by annealing the as-cast sample at 1223 K for 12 h and this alloy possessed great discharge capacity of 400 mAh g–1 [12]. But, the alloys with single 3R-A2B7 phase still have not been obtained. In order to comprehensively

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reveal the structural difference and electrochemical characteristics of these two type

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phase, a series of Re–Mg–Ni-based La0.59Nd0.14Mg0.27Ni3.3 alloy with certain proportions of 2H-A2B7 and 3R-A2B7 phases have been prepared by induction melting followed annealing at different temperatures. A crystallographic transformation model between 2H-A2B7 and 3R-A2B7 phase was proposed, and the electrochemical and kinetic characteristics of the La0.59Nd0.14Mg0.27Ni3.3 alloy were expounded in detail. 2. Experimental The La0.59Nd0.14Mg0.27Ni3.3 alloy was prepared by induction melting with the La,

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ACCEPTED MANUSCRIPT Nd, Mg and Ni constituent elements. The as-cast block La0.59Nd0.14Mg0.27Ni3.3 alloys were wrapped in a Ni foil shell, then annealed at 1148, 1173, 1223 or 1243 K for 24 h under an argon atmosphere and finally cooled to room temperature in a tube furnace.

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The heating rate was 4 K min–1 below 873 K and 1 K min–1 between 873 K and each target temperature. These four annealed alloys were named T1, T2, T3 and T4.

The alloy samples were crushed mechanically into particles of less than 400

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meshes for phase structure tests using X-ray diffraction (XRD) with a Rigaku D/Max

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2500/PC X-ray diffractometer (Cu–Kα radiation). The data collection profile was in the 2θ range from 10° to 80° with a step interval of 0.02° and a count time of 2 s per step. The Rietveld method was applied to analyze the collected data from the profile of the alloy using the Rietica software [13] and to calculate the phase abundance,

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lattice parameters and unit cell volumes for each sample.

Alloy electrodes were prepared by cold-pressing the mixture of carbonyl nickel powders and alloy powders in the weight ratio of 5:1 under 15 MPa. The sizes of the

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powders were between 200 and 400 meshes. The pellet was 10 mm in diameter. An

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electrochemical cell with a sintered Ni(OH)2/NiOOH as the positive electrode, a Hg/HgO electrode as the reference electrode and a 6 M KOH solution as the electrolyte was used to test the electrochemical measurements of the alloy electrodes at room temperature (298 K). For activation, charge/discharge cycling and electrochemical pressure–composition (P–C) tests, the experimental details have been reported previously [14]. For high-rate dischargeability (HRD) measurements, the electrodes were fully charged at a current density of 72 mA g–1 (0.2 C) and then

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ACCEPTED MANUSCRIPT discharged at current densities of 0.2 C, 1 C, 2 C, 3 C and 4 C. The electrochemical kinetic tests of the alloy electrodes, including linear polarization and potential-step measurements, were performed following the methods reported in our previous paper

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[15]. To study the pulverization and oxidation mechanism of the alloy samples during the charge/discharge cycles, XRD and SEM measurements were performed. For the

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cycling samples, a pellet was prepared by cold-pressing ca. 0.4 g of alloy powder

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(neither binders nor carbonyl nickel powders were added). Then, the pellet was wrapped in a nickel foam, welded around the nickel foam and measured in the half cell configuration. The alloy electrodes were charged at a current density of 72 mA g–1 for 10 h and then discharged at a current density of 72 mA g–1 until a cut-off voltage

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of 1.0 V was reached. The samples before cycling and cycling samples of 10, 40, 100 charge/discharge cycles were first collected, then washed with deionized water 4 times and dried under Ar.

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3. Results and discussion

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Fig. 1 shows the XRD profile of the T1–T4 alloys, and these alloys only contain the hexagonal Ce2Ni7-type (2H-A2B7) phase and the rhombohedral Gd2Co7-type (3R-A2B7) phase with different mass ratios. As Fig. 1 shows, with the annealing temperature increasing from 1148 K to 1248 K, all the diffraction peaks of the alloys become much sharper, which indicates an increase in the crystallinity of the alloy particles and a better crystal structure of the alloy. It is worth noting that the characteristic peaks of the 2H-A2B7 phase at approximately 33° and 37° become

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ACCEPTED MANUSCRIPT much more obvious with increasing temperature, especially at 1248 K. Fig. 2 presents the Rietveld refinement profiles of the T1, T2 and T4 alloys as examples. Phase abundance, lattice parameters and unit cell volumes are listed in

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Table 1 and are calculated from the data of Fig. 1 using the Rietveld method. Table 1 demonstrates that, at the annealing temperature of 1148 K, the phase composition of the T1 alloy is 72 wt.% for the 3R-A2B7 phase and 28 wt.% for the 2H-A2B7 phases.

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Moreover, as the temperature increases, the 3R-A2B7 phase abundance decreases to 63

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wt.% (T = 1173 K) and 35 wt.% (T = 1223 K), accompanied by an increase of the 2H-A2B7 phase abundance to 37 wt.% (T = 1173 K) and 65 wt.% (T = 1223 K), respectively. In particular, when the annealing temperature further increases to 1248 K, the 2H-A2B7 phase abundance surges to 97 wt.%. This crystal transformation from

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3R-A2B7 to 2H-A2B7 can be explained by the following reasons: one is that annealing treatment can make the phase composition of the alloy more homogeneous; the more stable phase tends to be generated during the annealing process, and the 3R-A2B7

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phase is mainly formed from the cooling process of molten metallic solution [10]. The

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other is that many crystal defects exist in the La2Ni7 alloy, but the 2H-type phase has fewer crystal defects than the 3R-type [16]. Therefore, the lower temperature is favorable for the formation of the 3R-A2B7 phase, and a higher temperature promotes the formation of the 2H-A2B7 phase. As a result, the 3R-A2B7 phase gradually transforms into the 2H-A2B7 phase with the elevation of the annealing temperature. In addition, although the single 2H-A2B7 phase has been prepared and studied [11,12,17], an alloy with a single 3R-A2B7 phase has not been obtained successfully.

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ACCEPTED MANUSCRIPT Therefore, the stable formation temperature for the 3R-A2B7 phase is not clear, and the exact transformation temperature between the 2H-A2B7 and 3R-A2B7 phases has not been determined. To explain the transformation mechanism clearly, the crystal

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transformation model from the 3R-A2B7 to the 2H-A2B7 phase is proposed using Diamond 3 software and is shown in Fig. 3 and was based on the identical chemical compositions of the two phases without requiring long-range atomic transport. The

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stacking layers of the 3R-A2B7 and 2H-A2B7 phase structures can be observed as the

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ABC- and AB-type blocks, respectively. The [ReNi5] slabs in block A beneath the 3R-A2B7 unit cell are slipped by – a/3 – 2b/3 to form the [ReNi5] slabs in block A below the 2H-A2B7 structure. Additionally, the [ReMgNi4] subunits in block A that connect the two [ReNi5] slabs have to be adjusted for accommodation; that is, the

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lower (Re+Mg) layer in Block A shifts together with the sliding [ReNi5] slabs, the upper (Re+Mg) layer slips by a/3– b/3 and the Ni layer between the two (Re+Mg) layers shifts by –2a/3 – 4b/3. By contrast, Block A on the 3R-A2B7 unit cell can be

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converted into Block B of the 2H-A2B7 structure with a slip by a/3 + 2b/3. The crystal

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transformation from 2H-A2B7 into 3R-A2B7 may follow the reverse process. P–C isotherms of the La0.59Nd0.14Mg0.27Ni3.3 alloys with different ratios of

2H-A2B7 and 3R-A2B7 phase compositions measured using electrochemical desorption are shown in Fig. 4, in which the following three features are observable: (1) For both the T1 alloy with the 3R-A2B7 phase as the main phase and the T4 alloy with the 2H-A2B7 phase as the main phase in P–C isotherms, there is only one discharge plateau region available (as shown in the illustration in Fig. 4): one is

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ACCEPTED MANUSCRIPT between 0.15 H/M and 0.85 H/M, where the hydrogen desorption plateau pressure is approximately 0.072 MPa. The other is between 0.16 and 0.86 H/M, and the hydrogen desorption plateau pressure is approximately 0.056 MPa. However, as a previous

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report states [18], the 2H-A2B7 phase and 3R-A2B7 phase usually show similar hydrogen desorption plateau pressures due to their similar stacking structures. Based on the results and combined with the study of Zhang et al. [19], we deduce that the

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3R-A2B7 phase has a higher plateau pressure upon hydrogen desorption than the

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2H-A2B7 phase. In addition, the T2 and T3 alloys exhibit two discharge plateau regions, which indicate two hydrogen desorption processes. For the T2 alloy with 63 wt.% 3R-A2B7 phase and 37 wt.% 2H-A2B7 phase, two different discharge plateau regions in the curves can be observed (as shown in the illustration in Fig. 4): one is between

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0.66 and 0.93 H/M and the other is between 0.37 and 0.58 H/M, with 0.079 MPa and 0.040 MPa as the equilibrium hydrogen pressures, respectively. As the 3R-A2B7 phase abundance decreases to 35 wt.%, accompanied by the increase of the 2H-A2B7 phase

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abundance to 65 wt.%, the discharge plateau regions of this alloy change; one is

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between 0.67 and 0.87 H/M and the other is between 0.32 and 0.6 H/M, and the hydrogen desorption plateau pressures are approximately 0.086 MPa and 0.043 MPa, respectively. Combining the above analysis and the same trend between the discharge plateau width and the phase abundance, we confirm that the higher equilibrium hydrogen pressure belongs to the 3R-A2B7 phase and the lower corresponds to the 2H-A2B7 phase of the T2 and T3 alloys. (2) Based on the analyses above, it can be determined that the T1 alloy possesses

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ACCEPTED MANUSCRIPT the highest equilibrium hydrogen pressure, which can be ascribed to the existence of the 72 wt.% 3R-A2B7 phase with a higher equilibrium hydrogen pressure. As the 3R-A2B7 phase transforms into the 2H-A2B7 phase, the equilibrium hydrogen pressure

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of the alloy decreases. In addition, for the T2–T4 alloy, the unit cell volumes of the alloy gradually shrink (as shown in Table 1) with the occurrence of the transformation, which results in more pressure being needed for the dehydrogenation of H atoms from

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the alloy bulk and in the elevation of the equilibrium hydrogen pressure.

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(3) The hydrogen desorption capacities of the T1–T4 La0.59Nd0.14Mg0.27Ni3.3 alloys are approximately 1.00 H/M (hydrogen atoms per metal atoms), 1.08 H/M, 1.06 H/M and 1.10 H/M, respectively, and the results are listed in Table 2. Except for the T1 alloy, these values are slightly higher than those of the hybrid La–Mg–Ni

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compound, La3MgNi14, with 1.03 H/M [16].

HRD curves of the La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes are shown in Fig. 5(a), and the discharge retention at the discharge current density of 1440 mA g–1 (HRD1440)

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is summarized in Table 3. The HRD can be defined with the following equation: (1)

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HRD = Cd / C72 × 100%

where Cd is the discharge capacity at the discharge current density Id and C72 is the discharge capacity at the discharge current density of 72 mA g–1. The T1 alloy with the 72 wt.% 3R-A2B7 and 28 wt.% 2H-A2B7 phases has the

lowest HRD, and the HRD1440 is 56.8%. As the 2H-A2B7 phase abundance increases to 65 wt.% and 97 wt.%, the HRD1440s of the alloy electrodes increases by 12.3% and 9.5%, up to 68.5% and 66.3%, respectively. It has been reported that a superior kinetic

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ACCEPTED MANUSCRIPT property of the alloy electrode is usually related to a higher hydrogen desorption plateau pressure [18,20]. As shown in Fig. 4, the tendency between the HRD and hydrogen desorption plateau pressure is almost the same, except for the T1 alloy. The

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above result implies that the 3R-A2B7 secondary phase is helpful for the HRD of the alloy with 2H-A2B7 as the main phase. However, the inherent characteristic of the phase will be the dominant factor in the performance of the high-rate dischargeability

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when the abundance is excessive.

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It is well known that the main factors influencing the HRD are the charge transfer reaction on the alloy surface and the hydrogen diffusion rate in the alloy bulk, which are characterized by the exchange current density I0 and the hydrogen diffusion coefficient D. Fig. 5(b) shows the linear polarization curves of the alloy electrodes at

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50% DOD. The polarization resistance Rp can be estimated from the slope of the curves, and the exchange current density I0 can be determined using the following formula:

(2)

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I0 = RTId / Fη (η→0)

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where R is the gas constant (J mol–1 K–1), T is the absolute temperature (K), Id is the applied current density (mA g–1), F is the Faraday constant (C mol–1) and η is the over-potential (mV), respectively. The calculated results are listed in Table 3. As shown in Table 3, the I0 of the T1 alloy with 28 wt.% 2H-A2B7 and 72 wt.% 3R-A2B7 phases is 329.51 mA g–1. When the 2H-A2B7 phase abundance increases to 65 wt.% and the 3R-A2B7 phase abundance decreases to 35 wt.%, the I0 of the T3 alloy increases to 395.03 mA g–1, but it slightly decreases to 370.24 mA g–1 when the

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ACCEPTED MANUSCRIPT 2H-A2B7 phase abundance of the T4 alloy is 97 wt.%. Fig. 5(c) shows the relationship between the exchange current density I0 and the HRD1440 of the alloy electrodes. A linear dependence of the HRD1440 on the exchange current density is established,

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which indicates that the charge transfer reaction on the alloy surface is the conclusive factor in the discharging process with high-rate current density.

The potential-step method is performed to determine the hydrogen diffusion rate

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in the alloy bulk. The fully charged electrodes are given a greater anodic potential so

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that the hydrogen diffusion rate becomes the rate-controlling step. The semilogarithmic curves of the anodic current density versus discharge time of the La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes are shown in Fig. 5(d). Assuming that a is 13 µm and the reaction time is sufficiently long, the hydrogen diffusion coefficient D can

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be calculated according to the slope of logi vs. t using the following equations: logi = log [ 6FD / da2(Co – Cs)] – π2Dt / 2.303a2

(3)

D = – 2.303a2 / π2 – d logi / dt

(4)

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where i, D, C0, Cs, d, a and t are the diffusion current density (A g–1), the hydrogen

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diffusion coefficient (cm2 s–1), the initial hydrogen concentration in the bulk of the alloy (mol cm–3), the hydrogen concentration on the surface of the alloy particles (mol cm–3), the density of the hydrogen storage alloy (g cm–3), the alloy particle radius (cm) and the discharge time (s), respectively. The hydrogen diffusion coefficient D values calculated using Equation (5) are summarized in Table 3. As shown in Table 3, the D value of the T1 alloy with a major 3R-A2B7 phase is 1.16×10–10 cm2 s–1, which is the highest. With a decreasing 3R-A2B7 phase, the D values decrease to varying degrees,

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ACCEPTED MANUSCRIPT indicating that the crystal transformation from the 3R-A2B7 phase to the 2H-A2B7 phase has a negative influence on the hydrogen diffusion rate in the alloy bulk. In addition, Table 3 shows that there is no distinct linear relationship between the

diffusion in the alloy bulk is not the rate-determining factor.

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hydrogen diffusion coefficient D values and the HRDs, implying that the hydrogen

The discharge capacities of the La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes within

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100 cycles are shown in Fig. 6, and the information from the curves is presented in

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Table 2. When the negative electrodes of Ni/MH batteries are charged for the first time, they experience the activation process. Hydrogen atoms are absorbed during the cycles until the maximum hydrogen storage capacity, in other words, the maximum discharge capacity, is obtained [21]. The fewer cycles needed for the batteries to reach

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the maximum discharge capacity, the better the activation property of the batteries. As shown in Fig. 6(a), the T2 and T3 alloy electrodes can be fully activated within 3 cycles, indicating that the alloy electrodes possess excellent activation properties.

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However, for the T1 alloy with the 72 wt.% 3R-A2B7 phase and T4 alloy with the 97

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wt.% 2H-A2B7 phase, the activation cycle number is five. Except for the T1 alloy, the evolution of the electrode discharge capacity essentially follows the same trend as that of the hydrogen storage capacity, which is calculated from the electrochemical P–C isotherms. The maximum discharge capacity (Cmax) of the alloy electrodes decreases from 386 mAh g–1 to 377 mAh g–1 as the 3R-A2B7 phase abundance decreases from 72 wt.% to 35 wt.%; this result is consistent with the study by Wang et al. [9]. However, the maximum discharge capacity of the nearly single 2H-A2B7 phase T4

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ACCEPTED MANUSCRIPT alloy electrode is 396 mAh g–1, which is superior to all the other alloys. As is established, the alloy component will be more homogeneous after the annealing treatment due to solid diffusion, and the degree of homogenization will increase with

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the elevation of the annealing temperature. In addition, the annealing treatment reduces the lattice strain and refines the grain sizes, which both facilitate compositional homogeneity. Therefore, the maximum discharge capacity of the T4

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alloy is superior to the other samples.

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The cyclic stability of the alloy electrodes within 100 cycles is also presented in Fig. 6(b). The capacity retention rate at the 100th charge/discharge cycle (S100) can be calculated using the following equation: S100 = C100 / Cmax × 100%

(5)

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where C100 is the discharge capacity at the 100th cycle and Cmax is the maximum discharge capacity during the cycles. The results are summarized in Table 2. The cyclic stability of the alloy electrodes increases from 88.5% to 92.4% as the 2H-A2B7

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phase abundance increases from 28 wt.% to 97 wt.%. Similar results have been

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reported by Gao et al. showing that the alloy with 93.4 wt.% 2H-A2B7 phase presented the best discharge capacity retention of 92.0% at the 100th cycle [22]. In addition, the discharge capacity at the 100th cycle was maintained at 366 mAh g–1 for the T4 alloy with the 97 wt.% 2H-A2B7 phase. To explore the mechanism underlying the improvement of the cyclic stability for the alloy, the microstructures of the T1–T4 alloy powders before and after cycling were compared, and the possible mechanisms are discussed below.

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ACCEPTED MANUSCRIPT Fig. 7 shows the XRD patterns of the fully discharged T1–T4 alloy electrodes at different cycles. As shown in Fig. 7(a) and (b), the diffraction peaks of the T1 alloy with 72 wt.% 3R-A2B7 phase and the T2 alloy with 62 wt.% 3R-A2B7 phase undergo a

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clear broadening and weakening in intensity after 40 cycles, indicating that the T1 and T2 alloys experience a certain degree of amorphous transition. By contrast, Fig. 7(c) shows that the T3 alloy with 65 wt.% 2H-A2B7 phase retains the main diffraction

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peaks, except for some low-intensity diffraction peaks that cannot be identified after

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more than 40 cycles. Moreover, after 10 cycles, diffraction peaks of La(OH)3 first appear, then diffraction peaks of Mg(OH)2 appear; they are caused by the oxidized La and Mg, respectively, on the surface of the alloy particles [23], meaning that La is easier to oxidize than Mg in the KOH electrolyte. In particular, there are no diffraction

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peaks corresponding to the oxides for the T4 alloy with 97 wt.% 2R-A2B7 phase, but the intensity of the diffraction peaks decreased, as shown in Fig. 7(d). The results above indicate that the 3R-A2B7 phase more easily becomes amorphous compared to

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the 2H-A2B7 phase. Therefore, the crystal transformation from the 3R-A2B7 phase to

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the 2H-A2B7 phase stabilizes the structure against amorphization, and then the crystal structure of the alloy can be effectively maintain after 100 charge/discharge cycles. A good crystal structure means that the alloy possesses a good reversibility of hydrogen absorption and desorption, less loss of the discharge capacity and further improves the cyclic stability. Fig. 8 presents the morphology of the corrosion products that developed on all of the alloys. The surfaces of the non-cycling alloys are smooth, but as soon as the

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ACCEPTED MANUSCRIPT corrosion process starts, the surfaces become rough. After 10 charge/discharge cycles, two morphologies are confirmed for the T1–T3 alloys. The T1 and T2 alloys show rough morphologies corresponding to corrosion products with irregular shapes. The

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T3 alloy shows needle-like and lamellar morphologies. For the T4 alloy, the surface of the alloy simply becomes rough and lacks any regularly shaped oxide products. Referring to the XRD patterns of the alloy electrodes after 10 charge/discharge cycles,

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the needle-like products may be associated with the La(OH)3 hydroxides. Meanwhile,

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for the T3 alloy, the diffraction peaks of La(OH)3 are more evident than those from the T1 and T2 alloys; thus, the morphology of the alloy is more defined. As the cycling time increases to 40 cycles, the alloy surface develops into a well-defined needle-like morphology. For the T1 and T2 alloys, the needles are quite large, up to the micron

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scale in size, while for the T3 alloy, the needles are smaller. When the cycling time extends to 100 cycles, both needle-like and lamellar products appear for the T1–T3 alloys; however, the change of surface morphology is not obvious for the T4 alloy.

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Therefore, the cyclic stability of the T4 alloy is superior to those of the other alloys.

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The lamellar products may correspond to Ni(OH)2 [24], but these products are not identified in the XRD analysis due to their low diffraction structure. 4. Conclusions

La0.59Nd0.14Mg0.27Ni3.3 alloys with specific proportions of the 3R-A2B7 and

2H-A2B7 phases have been obtained by induction melting followed by annealing at different temperatures. When the annealing temperature increases, the crystal transformation from the 3R-A2B7 phase to the 2H-A2B7 phase occurs, and the

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ACCEPTED MANUSCRIPT 3R-A2B7 phase decreases from 72 wt.% to 3 wt.% and is accompanied by the 2H-A2B7 phase increasing from 28 wt.% to 97 wt.%. Electrochemical P–C isotherms demonstrate that the 3R-A2B7 and 2H-A2B7 phases have quite different discharge

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plateaus, and the hydrogen desorption equilibrium pressure of the 3R-A2B7 phase is higher than that of the 2H-A2B7 phase. The 2H-A2B7 phase has better structural stability than the 3R-A2B7 phase. Therefore, the crystal transformation from the

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3R-A2B7 phase to the 2H-A2B7 phase heightens the structural stability of the alloy

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against amorphization and oxidation, further improving the cyclic stability. As a result, the discharge capacity retention at the 100th cycle of the alloy electrodes increases from 88.5% to 92.4%. Accompanied by the increasing 2H-A2B7 phase, the high-rate dischargeability of the alloy electrodes increases from 56.8% to 66.3%. The exchange

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current density has the same tendency as the HRD, indicating that the charge transfer reaction on the surface of the alloy electrode is the controlling step during the

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discharge process with large current densities.

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Acknowledgements

This work was financially supported by the National Natural Science Foundation

of China (NOs. 51571173 and 21303157), the Natural Science Foundation of Hebei Province (NO B2014203114) and Plan for Scientific Innovation Talent of Henan Province (NO 144100510009).

17

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Enhanced cycling stability and high rate dischargeability of (La,Mg)2Ni7-type hydrogen storage alloys with (La,Mg)5Ni19 minor phase, J. Power Sources 287 (2015) 237–246.

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ACCEPTED MANUSCRIPT Table 1 Lattice parameters and phase abundance of La0.59Nd0.14Mg0.27Ni3.3 alloy

T2 T3 T4

a (nm)

c (nm)

V (nm3)

2H-type

28

0.5026

2.4285

0.5312

0

3R-type

72

0.5027

3.6325

0.7948

0

2H-type

37

0.5026

2.4285

0.5312

0

3R-type

63

0.5027

3.6326

0.7948

0

2H-type

65

0.5026

2.4256

0.5308

0.075%

3R-type

35

0.5027

3.6285

0.7895

0.667%

2H-type

97

0.5028

2.4217

0.5302

0.188%

3R-type

3

0.4977

3.6413

0.7811

1.724%

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The unit cell volumes decreased of each phase in alloys compare to T1 alloy.

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a

Phase abundance (wt.%)

∆Va (%)

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T1

Phase Structure

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Sample

ACCEPTED MANUSCRIPT Table 2 Electrochemical performances of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes Nab

Cmaxc (mAh g–1)

C100d (mAh g–1)

S100e (%)

T1

1.00

5

386

339

88.5

T2

1.08

2

383

332

88.1

T3

1.06

3

377

341

90.5

T4

1.10

5

396

366

a

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Sample

Capacitya (H/M)

Hydrogen storage capacity from P–C isotherms The cycling numbers for the alloy electrodes to reach maximum discharge capacity c The maximum discharge capacity of the alloy electrodes d The discharge capacity at the 100th cycle of alloy electrodes e The capacity retention at the 100th cycle

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b

92.4

ACCEPTED MANUSCRIPT Table 3 Electrochemical kinetic properties of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes HRD1440 (%)

C1440 (mAh g–1)

Rp (mΩ)

I0 (mA g–1)

D (×10–10 cm2 s–1)

T1

56.8

226

77.97

329.51

1.61

T2

59.6

233

76.20

337.14

1.15

T3

68.5

268

65.04

395.03

1.19

T4

66.3

256

69.39

370.24

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Sample

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ACCEPTED MANUSCRIPT Figure Captions Fig. 1 (a) XRD profile of La0.59Nd0.14Mg0.27Ni3.3 alloys in the range of 10°–80°, (b) zoom of (a) in the range of 30°–50°.

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Fig. 2 Rietveld refinement profile of the alloy samples: (a) T1, (b) T2, (c) T4. Vertical bars below the profile show the positions of all possible reflection peaks: 2H-A2B7 and 3R-A2B7 phases.

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Fig. 3 Schematic illustration of crystal transformation from 3R-A2B7 phase to 2H-A2B7

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phase, the Re, (Re+Mg) and Ni atoms are represented as yellow, red/pink and blue spheres, respectively.

Fig. 4 Electrochemical desorption P–C isotherms of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes.

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Fig. 5 (a) HRD curves of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes, (b) Liner polarization curves of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes, (c) Correlation of the exchange current density I0 vs. the HRD1440 of the alloy electrodes, (d) Semilogarithmic curves of

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anodic current density vs. discharge time of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes.

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Fig. 6 Electrochemical discharge capacities of La0.59Nd0.14Mg0.27Ni3.3 alloy electrodes within 100 cycles.

Fig. 7 XRD profile at various discharge cycles of alloy samples, (a) T1, (b) T2, (c) T3, (d) T4.

Fig. 8 SEM images of alloy electrodes before cycling and after cycling, (a) T1, (b) T2, (c) T3, (d) T4.

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T1

30000

T2

25000

T3

20000

T4

15000 10000

Intensity (Count)

35000

(a)

5000

10

20

30

40

50

60

70

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0

80

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2θ (degree)

35000

T2

30000

T3

25000

T4

20000

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40

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45

2θ (degree)

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50

15000 10000 5000 0

Intensity (Count)

40000

(b)

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Rp = 3.61 Rwp = 4.97 S = 1.37

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(c)

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Fig. 2

Rp = 3.00 Rwp = 4.08 S = 1.36

Rp = 3.42 Rwp = 4.77 S = 1.39

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Fig. 3

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0.01 0.14 0.12 0.1 0.08 P/MPa

1E-3

T1

0.06 0.04

T2

T1 T2

1E-4

T3

T3 T4

0.02 0.2

0.3

0.4

1E-5

0.5

0.6

0.7

0.8

T4

0.9

H/M

0.2

0.4

0.6

H/M

1.0

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1.2

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0.1

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T1 T2 T3 T4

HRD (%)

60 -1

90 80 70 60 50

0

360

720

1080

90

30 0 -30 -60 -90

1440

-5

-4

-3

-2

-1

(d)

T3 390

Log (i (A g 1))

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350

1

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T1

-1.0

T2

-1.5

340

T2

T1

-2.0

58

60

62

64

66

68

-2.5

70

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HRD1440 (%)

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Fig. 5

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-1

I0 (mA g )

T4

360

320 56

1.0 0.5

380

330

0

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400

370

-1

T1 T2 T3 T4

Overpotential (mV)

Discharge current density (mA g ) (c)

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100

Current density (mAg )

(a)

0

500

1000

T4

T3

1500

2000

Time (s)

2500

3000

3500

ACCEPTED MANUSCRIPT (a) 400 -1

Discharge capacity (mAh g )

380 360 340 320

280 260 1

2

3

4

5

6

(b)

360

340

T1 T2 T3 T4

320

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40

60

Cycle number (n)

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Fig. 6

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9

10

11

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-1

Discharge capacity (mAh g )

380

300

7

Cycle number (n)

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0

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T1 T2 T3 T4

300

80

100

ACCEPTED MANUSCRIPT (b)

(a) • •

•

•

th

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•

Intensity (a.u.)

♦•

th

10 cycles •

•

• La(OH)3

♦

10

before cycling

30

40

50

60

70

• ♦

th

• La(OH)3 ♦ Mg(OH)2

20

before cycling

30

♦•

•

100 cycles th

•

40 cycles th

♦•

10 cycles

•

• La(OH)3 ♦ Mg(OH)2

40 50 2θ (degree)

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60

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Fig. 7

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40

60

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80

10

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30

th

100 cycles th

40 cycles

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Intensity (a.u.)

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80

•

•

Mg(OH)2

20

th

♦•

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th

10 cycles before cycling

40 50 2θ (degree)

60

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80

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Fig. 8

ACCEPTED MANUSCRIPT HIGHLIGHTS >Re–Mg–Ni-based alloys with specific proportion of 3R-A2B7 and 2H-A2B7 phases are obtained.

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>3R-A2B7 phase can transfer into 2H-A2B7 phase by increasing annealing temperature.

>3R-A2B7 phase is easier to become amorphous than 2H-A2B7 phase.

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>The 2H-A2B7 phase is favorable for the improvement of HRD and cyclic stability.