Abstract:

Degradation behaviors of La₂MgNi₉, La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ alloys were studied. The results indicate that severe pulverization and corrosion are important factors leading to the capacity deterioration. However, it is puzzled that corrosion of the electrochemical cycled alloys is aggravated, which is inconsistent with the result that La₂MgNi₉ present poor cycling stability andalso the assumption that alloy with high Mg content is easy to be corroded. Then, the intrinsic anti-corrosion and anti-pulverization characteristics were mainly focused in the first part of this work. Immersion experiments demonstrate that the Mg-rich phases are more easily to be corroded. The intrinsic anti-corrosion resistance of the three alloys presents an improved trend which is inversely proportional to the abundance of the Mg-rich phases. However, the intrinsic anti-pulverization ability just presents an inverse trend, which is closely related to mechanical property of the phase structures. LaNi₅ with the highest hardness is easy to crack, but the soft (La,Mg)Ni₂ is more resistant to crack formation and spreading, suggesting a possibility to improve the anti-pulverization ability by adjusting the phase constitution. In general, the weaker corroded extent of La₂MgNi₉ in the electrochemical test is attributed to its better intrinsic anti-pulverization capability though the intrinsic anti-corrosion of La₂MgNi₉ is worse. As to La₄MgNi₁₉ which possesses excellent intrinsic anti-corrosion resistant, enhancement of the anti-pulverization ability is the key issue to improve the cycling stability.

1 Introduction

Superlattice La-Mg-Ni based hydrogen storage alloys have received substantial attentions over the last decade because the excellent electrochemical performances used in nickel/metal hydride (Ni/MH) battery [1-5]. Up to now, A₂B₇ type alloys have been successful developed for the practical use [3]. However, AB₂ and AB₃ type alloys present poor cycling stability though the theoretical discharge capacities are higher than A₂B₇ type alloys [6-8]. In addition, A₅B₁₉ type alloys have been reported to possess good electrochemical performances, but they still need improvement to meet the practical application, especially on the cycling stability in the long-term reversible cycles [9-10].

It is well accepted that electrochemical capacity decrease of the metal hydride electrodes is caused by both the physical and chemical degradation [11-12]. In La-Mg-Ni system, factors affecting the capacity degradation were emphasized on pulverization and corrosion during the charge/discharge cyclings [13-17]. Corrosion leads to damage and disappearance of the phases which possess considerable hydrogen storage capacity. It has been reported that La-Mg-Ni alloys are easily to be corroded into La(OH)₃ and Mg(OH)₂ [13-15]. And these kinds of corrosion products are loose and passive which cannot protect the matrix for further corrosion [15-16]. Severe pulverization of La-Mg-Ni alloys during cycling had also been reported in many works [13-17]. Pulverization makes fresh surface of the electrodes alloys continuously exposed to the electrolyte and dramatically improves development of corrosion. Liu et al. classified the degradation process of the La-Mg-Ni-Co alloy into three stages: the pulverization and Mg oxidation stage, the Mg and La oxidation stage and the oxidation and passivation stage [14]. In addition, capacity degradation of the metal hydrides is closely related to the structural changes during absorption/desorption cycling. Our previous works demonstrated that transformation from crystallinity to amorphous viz. hydrogen induced amorphization (HIA) of La-Mg-Ni alloys occurred during the hydrogenation cycles and remarkably worsens both the gas-solid and electrochemical storage performances [18-19].

Understanding of the degradation mechanisms is the precondition for improvement of the cycling stability of the La-Mg-Ni based alloys. Several compounds including AB₂, AB₃, A₂B₇ and A₅B₁₉ type phase exist in this system, and the alloys usually present multi-phase microstructure. Though quiet a number of efforts have been applied on the degradation characters of the La-Mg-Ni based alloys, these works mainly focused on the overall capacity deterioration behaviors of the alloys. Diversity of the degradation characteristics of various compounds in this system is also lacking. In the present study, degradation mechanisms of three typical La-Mg-Ni alloys: La₂MgNi₉, La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ have been systematically investigated. In the first part of this work, corrosion and pulverization behaviors of the alloys, especially the intrinsic characteristics of the AB₃, A₂B₇ and A₅B₁₉ type La-Mg-Ni phases during absorption/desorption cycling were generated. In a following paper, HIA and its influence on the hydrogen storage properties are discussed.

2 Experimental materials and methods

2.1 Alloy preparation

The as-cast La₂MgNi₉, La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ alloy was prepared by induction levitation melting under argon atmosphere. The as-cast alloys were remelted twice for homogeneity. Appropriate excess of Mg was added in order to compensate for the evaporative loss of Mg during melting. Then the as-cast La₂MgNi₉, La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ alloys were annealed at 1143, 1173 and 1193 K respectively for 6 h protected in argon atmosphere.

2.2 Characterization

The sample was fine polished and then etched using a mixed etchant (including water, ethanol, acetic acid, picric acid, nitric acid and hydrochloric acid) at 343K. Then metallographic microstructure of the alloys was observed using a laser scanning confocal microscope (LSCM: Olympus-OLS4000). Phase constitution of the alloys was also characterized by a scanning electron microscopy (SEM: FEI-Qanta 400) under backscatter electron mode (BSE) applied on the unetched samples. The chemical composition of various phases was studied by energy dispersive spectroscopy (EDS) equipped in the SEM. Crystal structures of the alloys were measured by an X-ray diffractometer (XRD: Bruker-D8 Advance) with Cu Kα1 radiation. Micro-morphologies and selected area electron diffraction (SAED) were applied by means of a transmission electron microscopy (TEM: JEOL-2100 and FEI-F20) to examine the microstructural and crystallographic information. TEM samples were firstly crushed the bulk into fine powder, and then ultrasonic dispersion was performed in ethanol for 1800 s. Several drops of the mixed liquid were laid on a carbon membrane support on the copper grid, and dried in a vacuum oven.

Particle size of the cycled alloys was tested by a laser particle size analyzer (Malvern-Mastersizer 3000) where the alloy particles were dispersed by absolute alcohol. Oxygen content of the electrochemical cycled and immersed alloys was performed on a nitrogen/oxygen tester (NCS-ON3000). Before the oxygen test, samples were immersed in deionized water for 24 h, then washed using absolute alcohol twice to remove the residual KOH, and dried in a vacuum drying oven.

2.3 Hydrogen storage properties

Gas-solid cycling and P–C–T isotherm measurement were carried out by Suzuki -2SDWIN PCT system at 303K (Sievert’s type). Before the P–C–T analysis, sample was activated as follow: evacuated at 473 K for 2 h, placed to 303K, hydrogenated under 3Mpa H₂ (Purity 99.999%) pressure for 5 h, evacuated at 573 K for 2 h again. Each cycle consists of absorption at 2MPa for 600s and desorption by evacuating at 298K for 1200 s.

For the electrochemical measurement, the alloy particles (40-50µm) were mixed with carbonyl nickel powder in a weight ratio of 1:5 and cold pressed to form a pellet about 1g firstly. The pellets were then packed in a Ni foam substrate spot-welded with a Ni strip. The simulated three-electrode cell including a working electrode (metal hydride), a counter electrode (NiOOH/Ni(OH)₂) and a reference electrode (Hg/HgO) was installed. Before electrochemical test, the alloy metal hydride electrode was immersed in 6 M KOH aqueous solution for 1d. The measurement to get the maximum capacity and cycling stability was to charge at current density of 105 mA/g for 4h followed by a rest of 10min, then discharged at the same current density to the cut-off voltage of -0.6 V.

3 Results and discussions

3.1 Microstructure and hydrogen storage performances

LSCM and BSE micrographs of the three alloys are shown in Fig.1. Four contrasts can be detected in the La₂MgNi₉ alloy. The chemical quantitations of various contrast from EDS analysis are listed in Table 1, from which the four phases are speculated to be (La,Mg)Ni₂, (La,Mg)Ni₃, (La,Mg)₂Ni₇ and LaNi₅. Five crystal structures including CaCu₅-type, MgCu₄Sn-type, PuNi₃-type, Ce₂Ni₇-type and Gd₂Co₇-type are identified in XRD profile of the La₂MgNi₉ alloy, as shown in Fig.2. The structural parameters and phase abundance are refined and listed in Table 2. The results are in consistent with the metallographic observation that the main phase is (La,Mg)Ni₃, then (La,Mg)Ni₂and (La,Mg)₂Ni₇, but content of LaNi₅ is rare.

In case of the La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ alloy, metallographic and XRD characterization indicate that (La,Mg)Ni₂disappears, (La,Mg)₅Ni₁₉ emerges and LaNi₅ increases with elevation of the B-side stoichiometry. The main phase of the La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ alloy is (La,Mg)₂Ni₇ and (La,Mg)₅Ni₁₉ respectively, and the structural parameters and phase abundance are also listed in Table 2.

Fig.3 displays P-C-Tcurves of the alloys, and the detailed data are given in Table 3. Theoretically, hydrogenation capability increases with reduction of the B-side stoichiometry in the La-Mg-Ni based alloys. However, the maximum hydrogen absorption content of the La₂MgNi₉ alloy is slightly lower than the La_1.5Mg_0.5Ni₇ alloy. It is ascribed to the fact that some (La,Mg)Ni₂ which can hardly absorb and desorb hydrogen at room temperature [8], existing in the La₂MgNi₉ alloy. The three alloys have analogic hydrogen absorption plateau. But both the desorption pressure and the reversible hydrogen capacity elevate with increase of the B-side stoichiometry of the three alloys. Reversible hydrogen capacity of the AB₃-typed La₂MgNi₉ alloy is only 1.15 wt%, and the hysteresis effect is more evident than the other alloys.

Electrochemical discharge curves and performances of the alloys are shown in Fig.4 and Table 3 respectively. Discharge capacities of the La₂MgNi₉ and La₄MgNi₁₉ alloy are lower than the La_1.5Mg_0.5Ni₇ alloy. The lower discharge capacity of La₂MgNi₉ is due to the weak reversible hydrogen storage capacity. As to the La₄MgNi₁₉ alloy, it is attributed to high abundance of LaNi₅ which is unsuited for the electrochemical application without alloying [20]. Furthermore, La_1.5Mg_0.5Ni₇ presents better cycling stability than the other two alloys. Capacity retention after 100 cycles of the La₂MgNi₉ alloy is similar with that of the La₄MgNi₁₉ alloy.

3.2 Degradation characteristics after electrochemical cycling

From morphology and EDS results of the alloys, it is clear that pulverization and corrosion have occurred after electrochemical cycling by 100 times (only La₂MgNi₉ alloy presents in Fig.5). XRD analysis shows that La(OH)₃, Mg(OH)₂ and La₂O₃ appear in the cycled alloys, as displayed in Fig.6. Likewise, morphology and SAED analysis of TEM confirm existence of La(OH)₃ combined with La₂O₃ (stick-like), Mg(OH)₂ (needle-like) and MgO (particles), which are marked with 1, 2 and 3 respectively as illustrated in Fig.7. The results are in consistent with the other literature studied on the corrosion products of a La_1.5Mg_0.5Ni₇ alloy [21]. Detailed determinations of TEM are provided in the supplementary information (Fig.S1-S3). In addition, size and amount of La(OH)₃ and La₂O₃ are obvious than that of Mg(OH)₂ and MgO, indicating that corrosion of La is significant in the electrochemical environment. Mg(OH)₂ and MgO are close to the alloy surface but very loose. It agrees well with the previous works that corrosion products of Mg are gel-type and cannot form a solid protection layer for further corrosion [15-16].

Compared among the three alloys, it is noteworthy that corrosion productions of the La₂MgNi₉alloy are less than the other alloys (see in Fig.6). Identically, oxygen contents of the electrochemical cycled alloys follow the order that La₂MgNi₉ < La_1.5Mg_0.5Ni₇ < La₄MgNi₁₉, indicating that the corroded extent are aggravated (see in Fig.8). It is puzzled that the result is inconsistent with the electrochemical performances that La₂MgNi₉ possesses poor cycling stability. It also disagrees with the consideration that high Mg content is harmful to the corrosion resistance in La-Mg-Ni based alloys [22-24]. In order to comprehend this fact further, the intrinsic anti-corrosion resistance of the three alloys was investigated next.

3.3 The intrinsic anti-corrosion properties

To avoid impacts of pulverization on the corrosion behaviors, the alloy particles with the same diameter (around 40 μm) were immersed in KOH solution at 60 °C for 15 d. Then the morphology, phase structure and oxygen content were measured for characterization of the intrinsic corrosion behaviors. SEM micrographs and EDS analysis of the alloy particles illustrate that severe corrosion occurred after immersion, the typical results are shown in Fig.9 (only La₂MgNi₉ alloy particles are given here). Compared to the electrochemical cycled alloys, the stick-like products which have been confirmed as composite of La(OH)₃ and La₂O₃, are remarkable in the immersed samples which is due to aggravated corrosion at higher temperature.

XRD profiles identify that the corrosion products are mainly La(OH)₃, but La₂O₃ cannot be detected in the immersed alloys, as shown in Fig.10. Coincidently, SAED by TEM found that the stick-shaped phase is single-phase La(OH)₃, as shown in Fig.11. The result suggests that La₂O₃ transforms to La(OH)₃ during evolution of the corrosion process. Besides, Mg(OH)₂ and MgO are also found existing in the immersed samples, and their morphologies are same with that in the electrochemical cycled alloys. However, Mg(OH)₂can only be detected in La₂MgNi₉ from the identifications of XRD, indicating that corrosion of Mg is violent in La₂MgNi₉. Fig.12 is the oxygen contents of the immersed alloys, from which severity of corrosion of the three alloys are La₂MgNi₉ > La_1.5Mg_0.5Ni₇ > La₄MgNi₁₉.

To provide detailed information of the relationship between the corrosion behaviors and phase constitution, immersion test applied on the massive samples has also been studied (the condition is same with that of the powder samples). Fig.13 shows the SEM-BSE micrographs of the immersed samples (only La_1.5Mg_0.5Ni₇ alloy are present here). Obviously, the corroded extent is inhomogeneous which is considered to be caused by differences of the anti-corrosion capabilities of the various phases. EDS analysis on two regions with diverse corrosion grades (as marked with 1 and 2 in Fig.13) shows no Mg but less O existing in region 1. Whereas, more Mg and O are detected in region 2 with severe corroded extent than region 1. Likewise, EDS-mapping indicates that the region possessing more Mg presents richer O, as shown in Fig.14. Similar result is more evident in the as-cast alloys, which is attributed to the inhomogeneous chemical composition and microstructure of the as-cast alloy, details can be seen in the supplementary (Fig. S4 and S5).

The aforementioned results demonstrate that the Mg-rich phases are easy to be corroded in the alkaline solution. It has been well demonstrated that Mg solubility in La-Mg-Ni alloys follows the order that (La,Mg)Ni₂> (La,Mg)Ni₃ > (La,Mg)₂Ni₇ > (La,Mg)₅Ni₁₉ > LaNi₅ [25]. Thus the intrinsic anti-corrosion resistances of various phases in the La-Mg-Ni system are considered to be according with the inverse trend. This result is in agreement with several works where AB₂and AB₃ type La-Mg-Ni alloys have suffered serious corrosion after electrochemical experiments [8, 23-24]. The tendency is also exactly identical with that the anti-corrosion resistance is inversely proportional to the abundance of the Mg-rich phases. La₂MgNi₉ presents worse anti-corrosion capability because contents of the Mg-rich (La,Mg)Ni₂and(La,Mg)Ni₃ arehigher thanthe other two alloys. However, trend of the intrinsic anti-corrosion resistance is opposite to the corrosion extent of the three alloys after electrochemical cyclings. Concern to the fact that corrosion extent of the electrode alloys is also closely related to severity of pulverization during the electrochemical charge/discharge process, the pulverization properties of the alloys are carefully characterized then.

3.4 The intrinsic anti-pulverization properties

In order to avoid influence of the additives in the electrochemical test on characterization of the intrinsic pulverization behaviors, the alloys are gaseous hydrogenated and dehydrogenated for 30 cycles. Morphology observation indicates that remarkable pulverization has occurred where decrease of the particle size and emergence of cracks can be seen clearly in the cycled alloys, as shown in Fig.15 (only La₂MgNi₉ alloy are present here).

Then the particles sizes before (S_b) and after (S_a) cycling are measured and the size retention is calculated by S_b/S_a. It (Fig.16) shows that severity of pulverization for the three alloys are La₂MgNi₉ < La_1.5Mg_0.5Ni₇ < La₄MgNi₁₉, which is just contrary to the tendency of the corrosion extent after the immersion experiment. Combined with the results of the intrinsic anti-corrosion and pulverization characterization, we can conclude that the weaker corrosion extent of La₂MgNi₉ in the electrochemical test is attributed to its better intrinsic anti-pulverization capability though the intrinsic anti-corrosion of La₂MgNi₉ is worse.

It has been well accepted that pulverization is induced by the cell volume expansion upon hydrogen absorption [11-12]. Thus, large volume change leads to severe pulverization. Unfortunately, exact measurement of the volume expansion in the present work is difficult due to the multi-phase microstructure. Instead, we summarize the volume changes according to other experimental works where microstructures of these alloys are all sing-phase to ensure the accuracy as far as possible. Based on the data as listed in Table 4, there is no regular trend for the volume changes among the various structures in La-Mg-Ni system. And, no special relationship between the reported volume expansion data and the pulverization performances in the present work can be found. Besides, pulverization is believed to depend on the mechanical properties of the alloys [11-12]. Alloys with the more ductile character are more resistant to pulverization than the brittle materials. Usually, hydrogen storage alloys are hard and brittle, thus measuring ductility directly is difficult. Alternatively, Vickers hardness has been used to evaluate the preference of pulverization for the hydrogen storage alloys. And, previous works have found an almost inverse relationship between Vickers hardness and the pulverization rate [11-12, 30], suggesting the availability of Vickers hardness measurement on characterization of the anti-pulverization ability.

Fig.17 gives indentations of the various phases in La₂MgNi₉ and La_1.5Mg_0.5Ni₇ alloy. Evolution of Vickers hardness can be seen in Fig.18. It presents a linear relation between the Vickers hardness and B-side stoichiometry of the structures, which also agrees well with the pulverization behaviors of the alloys. Obviously, the mechanical property is an important factor affecting the anti-pulverization ability in the La-Mg-Ni phases. It is found that hardness differences of the AB₃, A₂B₇ and A₅B₁₉ type phase are small. Under low loading of the hardness test, micro-cracks can hardly be observed in all the above three phases. Since the test force increase, micro-cracks can be seen in all these phases, but there is no obvious difference between them. Differently, LaNi₅ is the hard phase, but hardness of (La,Mg)Ni₂ is far more lower than the other phases.

To comprehend more understanding on the crack formation of various phases, a massive sample with a polished surface was partial charged by electrochemical method, and the morphology and distribution of crack was observed. To highlight character of the hard and soft phase, the as-cast La₂MgNi₉ alloy was selected for the high abundance of LaNi₅ and (La,Mg)Ni₂. Microstructure characteristics of the as-cast La₂MgNi₉ alloy are given in the supplementary (Fig.S6-S7). As shown in Fig.19, quite a number of cracks can be observed in the sample which is only charged for 10 min. Most of the cracks exist in LaNi₅ with the darkest contrast in the BSE image. One reason is that LaNi₅ is the catalytic phase that primarily charged in the La-Mg-Ni system [31-32]. More importantly, it also ascribes to the brittle character of LaNi₅ which agrees well with the above result that the hard phase is easy to crack formation.

It is noteworthy that cracks are often stopped in front of (La,Mg)Ni₂. Obviously, the soft phase is more resistant to crack formation and able to prevent the crack spreading. Similar result has been reported in other literatures where ductile secondary phases are believed to be beneficial to the cycling stability [33]. According to the above results, we can conclude that La₄MgNi₁₉ alloy is easy to pulverization as the high abundance of the hard phases LaNi₅ and (La,Mg)₅Ni₁₉. As to La₂MgNi₉, little LaNi₅ but existence of the soft (La,Mg)Ni₂ and (La,Mg)Ni₃ make it more resistant to crack emergence. These findings enlighten a way to improve the anti-pulverization ability by introduction appropriate abundance and distribution of soft secondary phases.

4 Conclusions

In the present study, corrosion and pulverization behaviors of three typical La-Mg-Ni alloys: La₂MgNi₉, La_1.5Mg_0.5Ni₇ and La₄MgNi₁₉ have been systematically investigated. All the alloys present multi-phase microstructure with (La,Mg)Ni₃, (La,Mg)₂Ni₇and (La,Mg)₅Ni₁₉ as the main phase respectively. La_1.5Mg_0.5Ni₇ possesses better electrochemical properties among the three alloys. It is found that pulverization and corrosion with the main product La(OH)₃, combined with La₂O₃, Mg(OH)₂ and MgO, have occurred after the electrochemical cycling. The overall corrosion extent of the electrochemical cycled alloys follow the order that La₂MgNi₉ < La_1.5Mg_0.5Ni₇ < La₄MgNi₁₉. Immersion test demonstrate that the Mg-rich phases are easy to be corroded in the alkaline solution. The intrinsic anti-corrosion resistance are found to be La₂MgNi₉ < La_1.5Mg_0.5Ni₇ < La₄MgNi₁₉, which is inversely proportional to the abundance of the Mg-rich phases. However, the intrinsic anti-pulverization ability just presents an inverse trend that La₂MgNi₉ > La_1.5Mg_0.5Ni₇ > La₄MgNi₁₉. It is found that the mechanical property is an important factor affecting the anti-pulverization ability. Vickers hardness elevates with increase of the B-side stoichiometry of the various phases, which agrees well with the pulverization behaviors of the alloys. Furthermore, LaNi₅ with the highest hardness is found to be easy to crack formation, but the soft (La,Mg)Ni₂ is more resistant to crack formation and able to prevent the crack spreading. The weaker corrosion extent of La₂MgNi₉ in the electrochemical test is attributed to its better intrinsic anti-pulverization capability though the intrinsic anti-corrosion of La₂MgNi₉ is worse.

Acknowledgments

The authors are grateful to the Natural Science Foundation of China (NO. 51371094) and Natural Science Foundation Application of Inner Mongolia (NO.2014MS0526) for financial support.

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