Stages of mechanical alloying during the synthesis of Sn-containing AB5-based intermetallics

international journal of hydrogen energy 35 (2010) 6057–6062

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Stages of mechanical alloying during the synthesis of Sn-containing AB5-based intermetallics N.M. Cero´n-Hurtado a,b, M.R. Esquivel a,b,c,d,* a

Instituto Balseiro (UNCu and CNEA), Centro Ato´mico Bariloche, Av. Bustillo km 9.5, R8402AGP S.C. de Bariloche, Argentina Comisio´n Nacional de Energı´a Ato´mica, Centro Ato´mico Bariloche, Av. Bustillo km 9.5, R8402AGP S.C. de Bariloche, Argentina c Consejo Nacional de Investigaciones Cientı´ficas y Te´cnicas, Argentina d Centro Regional Universitario Bariloche, (UNCo), Quintral 1250, R8400FRF S.C. de Bariloche, Argentina b

article info

abstract

Article history:

The mechanical alloying of a La025Ce0.52Nd0.17Pr0.06–Ni–Sn mixture is studied by X-ray

Received 25 November 2009

diffraction, Scanning Electron Microscopy, Energy Dispersive Spectroscopy and Differential

Accepted 16 December 2009

Scanning Calorimetry. Four stages were identified and characterized. Initial stage was

Available online 13 January 2010

observed at integrated milling times (tm) between 0 and 30 h. It is dominated by fracture of

Keywords:

between 30 and 50 h. Both fracture and cold welding controls the process. At this stage,

AB5

compositional changes are detected due to solid–solid reaction. Ni and Sn particles are

the larger particles of Sn and La0.25Ce0.52Nd0.17Pr0.06. Intermediate stage is observed

Mechanical alloying

alloyed in the larger particles of La0.25Ce0.52Nd0.17Pr0.06. The Final stage is observed between

Hydrogen

50 and 70 h. At this stage, the cycle of fracture and cold welding reaches steady state and no further changes in chemical composition are observed. At completion stage, only refinement is observed. A La0.25Ce0.52Nd0.17Pr0.06Ni4.7Sn0.3 intermetallic is obtained. The compound is thermally stable up to 180  C in Air. ª 2009 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved.

1.

Introduction

The intermetallics based on AB5’s are widely used in different applications related to the hydrogen-based economy [1]. AB5’s are used in Nickel-hydride batteries and thermal compression of hydrogen (TCH) [2,3]. The versatility of the use of these intermetallics is based on the correlation between hydrogen sorption properties and composition [4]. In AB5’s, the A and B sites are occupied by lanthanides and transition metals, respectively. The most representative element in the B-site is Ni. The partial replacement of Ni by Sn affects the hydrogen equilibrium properties, the endurance to corrosion and cycling and the thermodynamic properties [5,6]. Although some of the effects of Sn in AB5

properties have been evaluated, research is still incomplete on the subject [6,7,8,9]. Most of these works were done using high temperature methods to synthesize Sn-containing AB5 [4,5,6,7,8,9]. Despite the fact that mechanical alloying presents various advantages over those methods including stoichiometry control and possibility of scaling up, little research was done on the synthesis of these compounds using this technique [10]. In this work, a Sn-containing AB5 intermetallic is synthesized by low energy mechanical alloying. The stages occurring during the processing of the Sn–Ni–La0.25Ce0.52Nd0.17Pr0.06 mixture are characterized and the governing processes are identified. The findings will be used to extend previous results obtained in an integral research program that encompasses

* Corresponding author. Comisio´n Nacional de Energı´a Ato´mica, Centro Ato´mico Bariloche, Av. Bustillo km 9.5, R8402AGP S.C. de Bariloche, Argentina. Tel.: þ54 02944 445156; fax: þ54 02944 445299. E-mail address: [email protected] (M.R. Esquivel). 0360-3199/$ – see front matter ª 2009 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2009.12.072

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synthesis, improvement and application to THC of hydride forming intermetallics [11,12].

2.

Experimental

Chunks of Sn (Berna, 99%), La0.25Ce0.52Nd0.17Pr0.06 [11] and powder Ni [11] were set with stainless steel balls in a stainless steel chamber under Ar atmosphere. The mixture was mechanically alloyed in an Uni-Ball-Mill II apparatus (Australian Scientific Instruments). A balls/sample mass ratio of 28.7 was selected. Handling and sampling of material during processing are reported elsewhere [11]. Room temperature X-ray diffraction was achieved on a Philips PW 1710/01 Instrument with Cu Ka radiation (graphite monocromator). Diffraction patterns were analyzed by the Rietveld method using DBWS software [13]. Particles size and morphology were observed by Scanning Electron Microscopy (SEM). Strain and crystallite size effects were estimated from diffraction peaks by assuming empirically a Gauss distribution and a Cauchy (Lorentz) component, respectively [14]. Chemical composition was verified by Energy Dispersive Spectroscopy (EDS) analysis. Thermal stability of the compound was analyzed by Differential Scanning Calorimetry (DSC 2970, TA Instruments). Measurements were done at 5  C min1 under Ar and Air.

3. 3.1.

Results and discussion Characterization of the reactants

The diffractograms of the starting La0.25Ce0.52Nd0.17Pr0.06, Ni and Sn are shown in Fig. 1a–c, respectively. The lanthanides alloy presents low crystallinity and wide peaks. Unlike this

component, both Sn and Ni present thin, well defined diffraction peaks. It is an indication of large crystallites and low strain values. These structures were refined using the Rietveld method [13]. Sn structure presents preferred orientation. It might be originated during the synthesis method since Sn is presented as folded sheets in its commercial form. Instead, Ni is presented as powder and no preferential orientation is observed. The structural parameters of both elements were refined using the Rietveld method. Cubic Ni was refined according to Fm3m space group with Ni distributed in Wyckoff positions 4a (0,0,0). Tetragonal Sn structure was refined according to I41/ amd space group with Sn distributed in Wyckoff positions 4a (0, 0, 0). An example of the refinements is shown for Ni in Fig. 1c. Experimental values are shown in black lines and refinement data in hollow dots. The bottom line in Fig. 1c is the difference between calculated and observed data. A summary of structural results is presented in Table 1. In these data, Rwp stands for the goodness of the fit. Sn cell parameters are smaller than the reference values as observed in Table 1. An SEM image of the initial Ni is presented in the upper right inset of Fig. 1. Flower-like particles with size distribution closer to 10 mm is observed. These sizes are smaller than the chunks of Sn and La0.25Ce0.52Nd0.17Pr0.06. It means that initial stage of milling will affect more the last particles before effective mechanical alloying takes place.

3.2.

Stages of milling

3.2.1.

Initial stage

The experimental patterns of Ni and Sn and the diffractogram of samples corresponding to tm values of 3 h, 15 h, 20 h and 30 h are shown in Fig. 2a–f, respectively. At 3 h, the peaks of Sn are still observed. Amorphization due to milling is produced and the peaks of this metal are not longer observed at 15 h. Lanthanide alloy chunks present low crystallinity and the peaks are not clearly observed except for the one at 2q ¼ 27.55 (dashed line). Ni peaks are clearly seen at these tm values. The estimation of the a parameter of Ni at different tm is presented in Table 1. Rietveld refinements of milled samples were done

Table 1 – Summary of Rietveld refinements results. For comparison, the values of the reference structures of both Ni and Sn are presented. In this table, Rwp stands for the goodness of the fit. Element

Cell parameters a  0.005 c  0.005 ˚) ˚) (A (A

Fig. 1 – XRD–SEM results a) Diffractogram of La0.25Ce0.52Nd0.17Pr0.06; b) Diffractogram of Sn; c) Diffractograms of Ni. Upper right inset. SEM image of initial Ni particles.

Ni Sn Ni JCPDF N 040850 Sn JCPDF N 010926 Ni milled (3 h) Ni milled (6 h) Ni milled (10 h) Ni milled (20 h) Ni milled (30 h)

a, b, g

V  0.5 ˚ 3) (A

Rwp (%)

3.517 5.787 3.523

– 3.535 –

90 90, 90, 120 90

43.50 105.62 43.72

10 10 –

5.819

3.175

90

107.52

–

3.521 3.521 3.523 3.510 3.509

– – – – –

90 90 90 90 90

43.76 43.76 43.76 43.24 43.20

10 9 11 15 15

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Fig. 2 – Diffractogram of reactants and as-milled samples. a) Ni; b) Sn; c) Milled 3 h; d) Milled 15 h; e) Milled 20 h; f) Milled 30 h.

with 2q ¼ 20–40 zone excluded. No appreciable changes are observed in this structural parameter. Therefore, no extended solubility between Ni and Sn in the Ni-rich side occurs in agreement with the corresponding binary phase diagram [15]. At 30 h, the formation of a new phase occurs on the (30–45 ) 2q range as a bump is observed along with the peaks of Ni and the peak of La0.25Ce0.52Nd0.17Pr0.06 (dashed line). It is clear from these results that chemical composition is not homogeneous at this stage. This feature is characteristic of this stage [16]. A mosaic of SEM images is presented in Fig. 3. Samples withdrawn at 3 h (a), 20 h (b) and 30 h (c) are depicted. The

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bigger particles of sizes close to 200 mm in Fig. 3a are chunks of La0.25Ce0.52Nd0.17Pr0.06 are fractured due to the impingements of the balls. Flat particles with sharp edges are the predominant shapes. This morphology is characteristic of a stage controlled by fracture [11,16]. At the end of the initial stage, the cycle of fracture and cold welding also affects the morphology (Fig. 2f and Fig. 3c). The formation of the new phase observed in Fig. 2f is in agreement with the change of the shape of the particles as shown in Fig. 3a–c. The predominant flat shape evolves to agglomerates of particles. At this stage, the energy of the milling was used to fracture the bigger ductile particles of La0.25Ce0.52Nd0.17Pr0.06. The smaller particles of Ni are less affected as noticed in the peaks of Fig. 2c–f. The microstructural parameters of Ni were calculated for some of these peaks. The summary of results is presented in Table 2. As observed, the strain (s) values are almost the same. But a diminution of one order of magnitude is observed in the crystallite size (d ). Because of the smaller size and less ductility than the lanthanide alloy and Sn, Ni is less affected at this stage. Then, the formation of the new phase is due to solid–solid reaction of the smaller particles of Ni within the bigger particles of La0.25Ce0.52Nd0.17Pr0.06 and Sn which serve as the matrix for the formation of the new intermetallic. It is confirmed by the EDS measurements shown at these tm in Table 3. Results are expressed as atomic percent. Some of the particles at 3 h show no traces of Sn. The particles withdrawn at 10 h and 30 h show the presence of Sn. From these results, it can be inferred alloying and therefore, the formation of a new phase.

3.2.2.

Intermediate stage

The formation of a new phase as observed in Fig. 2f. It implies a shift in the governing process that controls the mechanical alloying. The process of fracture does not produce chemical modifications which are only due to cold welding [11,16].

Fig. 3 – SEM images at different mechanical alloying stages. The mosaic also displays the governing processes.

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Table 2 – Summary of microstructural calculations on Ni peaks. Rietveld refinements were made with 2q 20–40 excluded zone. Rwp stands for the goodness of the fit. Sample/treatment

hkl

˚ )  10 d (A

s (%)  1

Initial Ni

111 200

1343 1141

1 1

Milled 10 h

111 200

554 442

3 3

Milled 20 h

111 200

490 370

3 4

Milled 30 h

111 200

405 280

4 5

Therefore, the mechanical alloying evolves from initial to intermediate stage as the main control of the process changes from fracture to fracture-cold welding shared control. This shift is also observed in the morphology of the samples. Flat particles with sharp edges predominant morphology evolves to spheroid shape like as observed in Fig. 3d. An increasing control of cold welding is correlated to the decrement of the size of the bigger particles of La0.25Ce0.52Nd0.17Pr0.06 and Sn. It also indicates that closer sizes are reached between the constituents and solid–solid reaction is achieved. At this stage, chemical composition is still not homogeneous [11,16]. This assessment is verified by the analysis of the reaction products presented in Table 3 and Fig. 2f. Starting Ni is presented for comparison in Fig. 4a (dashed lines). The formation of a new phase along with the peaks of Ni is observed in the 25–45 2q range of Fig. 4b and 30–45 2q of Fig. 2f.

3.2.3.

Final and completion stages

The SEM image of a sample extracted at 70 h is shown in Fig. 3e. The corresponding diffractogram is presented in Fig. 4c. As mechanical alloying progresses, the cycle of fracture and cold welding reaches equilibrium because the highly reactive surfaces created by fracture are welded together by reactive particles. This behavior forces the particle shape to be spheroid-like as observed in the image of Fig. 3e. Since this processing affects the bulk of the sample and no bigger

Table 3 – Summary of EDS measurements for samples extracted at final stage (70 h). Calculation was done by distributing Lanthanides in A-site and Ni and Sn in B-site. Sample/treatment Particle size m  m

Element (At %  2) La Ce Nd Pr Ni Sn

As-milled 3 h As-milled 10 h As-milled 30 h

Average Average Average

19 43 4 10 5 9

As-milled 100 h

15  15 15  15 10  10 10  10 10  10 10  10

25 26 25 26 25 25

52 51 52 52 52 52

5 12 18 – 1 3 67 5 1 3 68 4 17 17 16 17 17 17

5 6 7 5 6 6

94 95 94 95 94 94

6 5 6 5 6 6

Fig. 4 – Diffractograms of Ni and as-milled samples. a) Initial Ni; b) Milled 50 h; c) Milled 70 h; d) Milled 100 h.

particles remain, the chemical composition reach its final value and the formation of the intermetallic is obtained as presented in Fig. 4c. In this figure, the characteristic family planes (101), (110) and (111) of an AB5-based structure are indicated [16]. The figure also shows the presence of Ni peaks (dashed lines). Both structures constitute the coexisting ones on any rich Ni side, AB5-Ni based phase diagram [17]. If the milling process is continued, the completion stage is reached [11,16]. The particle size is refined and the surfaces become highly reactive. It enhances the formation of agglomerates and an increasing control of cold welding occurs in the process [16] leading to the successive formation of bigger particles. An SEM image of a sample extracted at this stage at 100 h is shown in Fig. 3f. It is seen large agglomerates along with the formation of cold welded particles (indicated with white circles). Since no chemical composition changes are possible at this stage, the process shifts from mechanical alloying to mechanical milling and progressive amorphization of the phases present might be observed. This behavior is observed in the diffractogram of Fig. 4c,d. The composition is not changed since the AB5 was already obtained. Ni is amorphized and the crystallite size of the AB5 is increased as observed in the labeled peaks of Fig. 4c,d. The final composition was determined by EDS. Results are summarized in Table 3. Lanthanides are grouped in A-site as atomic %. Sn and Ni are grouped in the B-site as atomic %. The final composition obtained is La0.25Ce0.52Nd0.17Pr0.06Ni4.70Sn0.3.

3.3.

Thermal stability of the as-milled intermetallic

The thermal stability of the as-milled intermetallic was evaluated under Ar and Air from room temperature (TA) to 200  C at 5  C/min. It is the range of interest because THC’s based on AB5’s are usually operated between 10 and 90  C. The DSC curves are shown in Fig. 5. Under Ar atmosphere, no significant thermal evolution is observed and the curve remains

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evolves from sharp-edged particles to rounded ones as studied by SEM. The evolution of a rich-La intermetallic is observed by XRD and EDS indicating a change in chemical composition. The final stage occurs between 50 and 70 h. The process reaches steady state and composition its final value. At longer tm values, the particles are refined. Processing produces highly distorted particles surface. This fact favors the cold welding among them and the process shifts progressively from mechanical alloying to mechanical milling. The thermal stability of the as-milled samples was evaluated under Ar and Air. No phase transformation or disproportionation of the intermetallic was observed under Ar up to 200  C. The intermetallic resists thermal heating under Air up to 180  C. These results indicate that the obtained intermetallic is structurally stable under inert and aggressive atmosphere at temperatures higher than the operating range. The hydrogen sorption properties of these intermetallics are currently being evaluated. These results constitute the subject of an incoming paper. Fig. 5 – DSC curves. Baselines are in dotted lines a) Asmilled intermetallic under Ar. b) As-milled intermetallic under Air c) La0.25Ce0.52Nd0.17Pr0.06 under Air.

closer to the baseline (Fig. 5a). No phase transformation, reordering or disproportionation of the initial structure is observed. This behavior was already reported in pseudo binary AB5’s obtained using the same MA device and synthesis conditions [12]. Under Air atmosphere, no thermal event was detected up to 180  C (Fig. 5b). At this temperature, the on-set of the thermal decomposition of the intermetallic occurs leading to the formation of La0.25Ce0.52Nd0.17Pr0.06O2, NiO and SnO. For comparison, the DSC curve of La0.25Ce0.52Nd0.17Pr0.06 is presented (Fig. 5c). The exothermic on-set at 78  C corresponds to the starting temperature of the oxidation of La0.25Ce0.52Nd0.17Pr0.06. At these temperatures, Ni and Sn do not react with oxygen [18]. Therefore, no free La0.25Ce0.52Nd0.17Pr0.06 remains in the as-milled intermetallic and no disproportionation of La0.25Ce0.52Nd0.17Pr0.06Ni4.7Sn0.3 occurs up to 180  C. The result is important because it indicates that the bulk intermetallic is stable under inert and aggressive atmospheres at temperatures at least 90  C higher than the usual operation temperature range.

4.

Conclusions

In this work, A La0.25Ce0.52Nd0.17Pr0.06Ni4.7Sn0.3 intermetallic was synthesized by low energy mechanical alloying. The stages occurring were identified and characterized. Four stages were observed. Initial stage occurs from 0 to 30 h. No significant changes were observed in the microstructural parameters of Ni as presented in Table 1. Then, the stage is dominated by the fracture and amorphization of the larger particles of La0.25Ce0.52Nd0.17Pr0.06 and Sn as observed in Figs. 1 and 2a. At tm values between 30 and 50 h, Intermediate stage is observed. The still larger particles of La0.25Ce0.52Nd0.17Pr0.06 became the matrix where Ni and Sn are welded. At this stage, cold welding affects the process since particles morphology

Acknowledgements The authors wish to thank to the Agencia Nacional de Promocio´n Cientı´fica y Tecnolo´gica of Argentina (Project PICT 33473 and Project PAE-PICT 00158) and to Universidad Nacional de Cuyo (Project 06/C310) for partial financial support.

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