Crystallite sizes of LiH before and after ball milling and thermal exposure

Journal of Alloys and Compounds 454 (2008) 297–305

Crystallite sizes of LiH before and after ball milling and thermal exposure Angel L. Ortiz a , William Osborn b , Tippawan Markmaitree b , Leon L. Shaw b,∗ a

Departamento de Electr´onica e Ingenier´ıa Electromec´anica, Universidad de Extremadura, Badajoz, Spain b Department of Chemical, Materials and Biomolecular Engineering, University of Connecticut, 97 North Eagleville Road, U-3136 Storrs, CT 06269-3136, United States Received 4 August 2006; accepted 10 December 2006 Available online 16 January 2007

Abstract The powder characteristics of lithium hydride (LiH) as a function of high-energy ball milling condition are systematically investigated via quantitative X-ray diffraction (XRD) analysis. The results obtained from the XRD analysis are compared with those attained from scanning electron microscopy (SEM), transmission electron microscopy (TEM), and specific surface area (SSA) analyses. The thermal stability of the ball-milled LiH is also investigated in order to provide physical insights into its cyclic stability in hydrogen sorption and desorption cycles. The results indicate that ball milling is effective in obtaining nano-crystalline LiH powder which is relatively stable with retention of nano-crystals after thermal exposure at 285 ◦ C (equivalent to 0.58Tm ) for 1 h. The good thermal stability observed is attributed to the presence of many pores in the agglomerates at the ball-milled condition. These pores effectively prevent crystal growth during the thermal exposure. © 2006 Elsevier B.V. All rights reserved. Keywords: Metal hydrides; High-energy ball milling; Thermal stability; Hydrogen storage materials

1. Introduction High-energy ball milling can provide mechanical activation to solids. The mechanical activation obtained can increase the rates of adsorption, absorption, and chemical reactions [1–4]. Solution of gases in solids can also be enhanced by mechanical activation [3,4]. As such, high-energy ball milling has been widely used to increase the hydrogen sorption and desorption rates [5–13], to decrease the peak temperature for hydrogen uptake and release [6,14], and in some cases to increase the hydrogen storage capacity [15–19]. In spite of the paramount importance of mechanical activation, systematic investigation of the effect of high-energy ball milling on the powder characteristics of hydrogen storage materials is very limited. This is especially true for the Li–N–H and Li–Mg–N–H systems which have recently been investigated extensively owing to their potentials in the reversible hydrogen storage for fuel cell technologies [20–31]. There is no doubt that studies in this area are necessary in order to fully utilize high-energy ball milling to develop

∗

Corresponding author. Tel.: +1 860 486 2592; fax: +1 860 486 4745. E-mail address: [email protected] (L.L. Shaw).

0925-8388/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2006.12.035

viable hydrogen storage materials based on the Li–N–H and Li–Mg–N–H systems. Lithium amide (LiNH2 ), lithium imide (Li2 NH), lithium hydride (LiH), magnesium amide (Mg(NH2 )2 ), magnesium imide (MgNH), and magnesium hydride (MgH2 ) involved in the Li–N–H and Li–Mg–N–H systems are all brittle materials. Based on this, it is expected that the particle refinement of these amides, imides and hydrides during high-energy ball milling would be attained mainly through fracture with little cold welding. However, the melting temperatures of these amides, imides, and hydrides are relatively low if one considers that the local temperature at the collision site at the moment of collision can be as high as 500 ◦ C during ball milling [32,33]. Given such a thermal exposure, although very short, it would be expected that cold welding could proceed extensively during ball milling of these amides, imides, and hydrides. Clearly, a priori prediction would be difficult since the compound brittleness and their low melting temperatures have opposite effects. Whether the powder characteristics obtained via high-energy ball milling are stable or not in hydrogen sorption and desorption cycles is another important aspect of the mechanically activated hydrogen storage materials. The cyclic stability is an important issue because mechanically activated materials are typically in a

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high-energy state with high defect concentrations and, in many cases, with metastable structures [34–36]. Thus, mechanically activated materials have tendency to transform into thermodynamically stable states during hydrogen sorption and desorption cycles, and thus may gradually change their hydrogen sorption and desorption behavior. Lithium and magnesium amides, imides, and hydrides are especially prone to this potential problem because of their low melting temperatures. The present study is the first attempt to systematically investigate the powder characteristics of LiNH2 , Li2 NH, and LiH as a function of high-energy ball milling conditions. However, only LiH is studied in this first attempt because of the complexity in analyzing air-sensitive nano-particles. LiH is subjected to high-energy ball milling for different times, and its powder characteristics, including the crystallite size, size distributions, internal strains, and lattice parameters, are quantified as a function of the ball milling time via detailed X-ray diffraction (XRD) analysis. The results obtained from the XRD analysis are compared with those attained from scanning electron microscopy (SEM), transmission electron microscopy (TEM), and specific surface area (SSA) analyses. The thermal stability of the ballmilled LiH is also investigated in order to provide physical insights into its cyclic stability in hydrogen sorption and desorption cycles. 2. Experimental procedure 2.1. Material preparation LiH with 95% purity was purchased from Sigma–Aldrich. This as-purchased LiH powder is composed of single crystal particles, and has an average particle size of ∼1 ␮m with a size distribution ranging from 0.15 to 4 ␮m (see Fig. 1). The high-energy ball milling was conducted under ultrahigh purity argon (99.999% pure) using a modified Szegvari attritor, which has been shown to be effective in preventing the formation of the dead zone and producing uniform milling products within the powder charge [37]. The use of the ultrahigh purity argon atmosphere is to prevent oxidation of LiH during the ball milling process, as has been shown recently elsewhere [38]. The canister of the attritor and balls 6.4 mm in diameter were both made of stainless steels. The loading of balls and the powder to the canister was performed in a glove box filled with ultrahigh

purity argon, again to prevent oxidation. The ball-to-powder weight ratio was 60:1, the milling speed was 600 rpm, and the milling temperature was maintained at 20 ◦ C, achieved by water-cooling at a flowing rate of 770 ml/min. The ball milling time was a variable in this study to impart different degrees of milling intensity to the LiH powder, and ranged from 45 to 180 min. To study the thermal stability of the mechanically activated LiH, the powder ball-milled for 3 h was exposed to 285 ◦ C for 1 h in a pressure vessel under 1 atm of the ultrahigh purity argon. 285 ◦ C was chosen because this was the temperature at which the LiNH2 and LiH mixture exhibited 1 atm of hydrogen pressure [20]. Thermal exposure to 285 ◦ C would be necessary if 1 atm of hydrogen pressure is to be utilized for hydrogen sorption and desorption operation. This situation, of course, could change in the future if advancements are made in reducing the temperature at which 1 atm of hydrogen pressure is present. The ball-milled samples with and without the thermal exposure were all subjected to characterizations using XRD, SEM, and TEM, as described below.

2.2. X-ray diffractometry analysis The ball-milled LiH powders were all analyzed using XRD to determine the lattice parameters, the crystallite sizes, and the internal stresses. The aspurchased LiH powder was also analyzed, and utilized as the standard for elimination of the instrumental broadening from the LiH peaks in the ball-milled powders. Before being installed in the XRD holder, all LiH powders were mixed with pure silicon (Si) powder, which was used as the internal standard for the correction of the instrumental shifts in the position of the LiH peaks during the lattice parameter determination. For the XRD analysis, the 1 1 1 peak of both Si and LiH were collected via conventional powder diffractometry, with the powder mixture held inside a sealed capillary tube to prevent oxidation of LiH powder during the XRD data collection. The mixing of the Si and LiH powders and the subsequent loading of the powder mixture to the capillary tubes were both performed in a glove box filled with the ultrahigh purity argon. The capillary quartz tube had a wall of 0.01 mm thick and thus was transparent to the X-ray beam. The operation conditions for the data collection were Cu K␣ ˚ 40 kV, 40 mA, and 60 s count time per step using a D8 radiation (λ = 1.54183 A), ADVANCE diffractometer. The step sizes for the Si 1 1 1 and LiH 1 1 1 peaks were 0.02◦ and 0.01◦ step−1 , respectively. After XRD data collection, Voigt (V) and pseudo-Voigt (pV) functions with Cu K␣1 and Cu K␣2 components in the forms:



+∞

V (2θ) = I0,L I0,G





 2θ − 2θ − 2θ 2 −1 0

1+ π

× exp −π

 2θ − 2θ 2 

  pV (2θ) = I0 η 1 +

0

βG

d(2θ)

(1)

 2θ − 2θ 2 −1 0

w

 +(1 − η) exp −ln2

Fig. 1. Scanning electron microscopy image of the as-purchased LiH powder.

βL

−∞

 2θ − 2θ 2  0

w

 (2)

were fitted individually to each Si and LiH peak with the background being modeled by a first-order polynomial function. The Cu K␣2 component was assumed to have the same shape as the Cu K␣1 component, but with the half of its intensity and shifted toward higher angles according to the Bragg law (θ0,K␣ = arcsin(λK␣ sin θ0,K␣ /λK␣1 )). The fits were conducted iteratively by the Levenberg–Marquardt non-linear least-squares algorithm [39], and halted after complete convergence (χr2 < 1.5). In Eq. (1), I0 ,L and βL are the height and integral breadth of the Lorentz component of the Voigt function, respectively, and I0 ,G and βG are those of the Gauss component. In Eq. (2), I0 , w and η are the height, half width at half maximum and mixing parameter of the pseudo-Voigt function, respectively. Finally, θ and θ 0 in Eqs. (1) and (2) are the diffraction and Bragg angles, respectively. The lattice parameter of LiH in the as-purchased and ball-milled powders was determined from the position of their 1 1 1 peaks (θ 0 ,LiH ) using the Bragg law coupled with the plane-spacing equation for the cubic crystal system [40],

A.L. Ortiz et al. / Journal of Alloys and Compounds 454 (2008) 297–305 that is, through the following relation: √ 3λ a= 2 sin(θ0,LiH − θSi )

(3)

where λ is the radiation wavelength and θ Si is the instrumental shift measured for the 1 1 1 peak of the Si standard. θ Si was evaluated by subtracting the position of the 1 1 1 Si peak given in the 27-1402 reference PDF card from the one measured experimentally.The surface-weighted average crystal size (DS ) and the root-mean-square (eRMS ) microstrain of LiH in the ball-milled powders were calculated by the variance method through the following relations [41]: Ds =

108.81λ π(βL,h − βL,g ) cos(θ0,LiH − θSi )



π

eRMS =

(4)

360(2)1/2 tan(θ0,LiH − θSi )

2 2 βG,h − βG,g −

2(βL,h − βL,g )2 π

1/2 (5)

where the different parameters have been defined above. The subscripts h and g refer to the ball-milled LiH and as-purchased LiH, respectively. The crystallite size distribution (P(D)) of LiH in the ball-milled powders was calculated by an analytical model based on the line-broadening theory of XRD through the following relation [42]:



1

exp P(D) = √ 2πσD



1 D − 2 ln Dmed 2σ

2

(6)

with σ and Dmed being the median and the deviation of the lognormal crystallite size distribution, and are given, respectively, by:



σ=

− ln

Dmed =

(2.114 + 0.694ηf )2 (1.344 + 11.488ηf )

(7)

180λ √ π2 wf cos(θ0,LiH − θSi ) 0.084 + 0.718ηf



× exp

3 ln

(2.114 + 0.694ηf )2 (1.344 + 11.488ηf )



(8)

where the different parameters have the same meaning as defined above. The subscript f refers to the ball-milled LiH excluded the instrumental broadening, which was done in this study by a modified Stokes method [42]. In practice, a, DS , eRMS , and P(D) were calculated via introducing into the above equations the values of different parameters associated only to the Cu ˚ ). K␣1 component (λCu K␣1 = 1.54056 A

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SEM (FESEM) operating at 5.0 kV and 10−4 Torr. A 20 s sputter coating of gold–palladium deposited using an E5100 SEM coating unit at 2.2 kV and 20 mA was required to prevent charging on the SEM samples. The TEM characterization was completed using a JEOL 2010 FasTEM with a 200 kV thermionic source, and both bright- and dark-field image techniques were used.

3. Results and discussion 3.1. XRD analysis of the ball milling time effect Shown in Fig. 2 are selected 1 1 1 peaks of LiH with and without ball milling. The following two interesting features are noted after ball milling: (1) the peaks do not shift and (2) the peaks are increasingly broader with increasing ball milling time. The absence of peak shifting after ball milling excludes the existence of long-range internal stresses and/or macroscopic residual stresses acting on the LiH crystals [44–46], whereas the peak broadening could be, in principle, the consequence of: (i) crystallite size refinement, (ii) introduction of lattice defects (point, line, or planar defects), or (iii) both [44–46]. Shown in Fig. 3 is the lattice parameter of LiH before and after ball milling, determined with the aid of Eq. (3). It can be seen that the as-purchased LiH and the three ball-milled LiH have the same lattice parameter within the experimental error, and it is thus concluded that the ball milling process has not altered the crystal structure of LiH and its lattice parameter. With respect to the actual cause(s) of the peak broadening after ball milling, the profile fitting can provide invaluable insights through the shape parameters of the 1 1 1 XRD peaks. With increasing ball milling time, it is observed that the integral breadth of the Lorentz component (βL ) increases from 0.0004◦ for the as-purchased LiH to 0.2132◦ for the 180 min ball-milled LiH, whereas, on the contrary, the integral breadth of the Gauss component (βG ) remains the same at the as-purchased condition of 0.46◦ . Clearly, the unchanged βG is incompatible with the introduction of lattice defects within the LiH crystals

2.3. Specific surface area, SEM, and TEM analyses The specific surface area of the powder before and after high-energy ball milling was determined through nitrogen adsorption at 77 K based on the Brunauer–Emmett–Teller (BET) method [43] using a gas sorption analyzer (NOVA 1000). The loading of the sample (∼0.40 g) into a sample cell with a Teflon stem filler was performed in a glove-box filled with Ar of 99.999% purity. The measurement was performed immediately after the sample was loaded in the instrument without a degassing treatment. The relative pressure (P/Po ) was 0.05–0.3 and the reported SSA data were calculated based on five points BET method. The particle morphology of the powders before and after ball milling and thermal exposure was examined using SEM and TEM. Samples for electron microscopy were prepared by suspending 5 mg of LiH powder in 7 ml of 99.8% cyclohexane (Fisher C556-4) and ultrasonicating for 3 h. Gold–palladium coated glass coverslips and carbon film TEM grids, for SEM and TEM samples, respectively, were maintained at 90 ◦ C on a hotplate while a drop of the suspension was placed. The samples were allowed to evaporate at 90 ◦ C for 10 min before being removed from the hotplate. In all cases, the sample preparation was performed in an argon (99.999%) glove box. SEM characterization was performed using a JEOL 6335F field emission

Fig. 2. 1 1 1 XRD peaks for the as-purchased LiH (dotted line) and the 3 h ballmilled LiH (solid line) after correction by instrumental shifts. For comparison, the two peaks have been normalized by imposing the same maximum intensity. The XRD patterns of the 45 and 90 min ball-milled LiH are between those of the as-purchased LiH and 180 min ball-milled LiH, and are there omitted for the sake of clarity.

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Fig. 3. The lattice parameter of LiH as a function of ball milling time (close dots) and thermal exposure (open dot). The solid line is a guide for the eyes.

during ball milling, since the attendant internal stresses would induce additional Gauss broadening [41]. On the other hand, the increased βL is entirely consistent with the crystallite size refinement from the micrometer range to the nanometer range, since the presence of LiH nano-crystals would induce additional Lorentz broadening [41]. To confirm that the cause of the peak broadening in the ballmilled LiH is crystallite size refinement, not introduction of lattice defects within the crystals, the average crystallite sizes and internal stresses from the variance of the 1 1 1 XRD peaks have been evaluated. Shown in Fig. 4 are the crystallite sizes determined with the aid of Eq. (4). It can be seen that the crystallites in the ball-milled LiH powders are nanometer-sized, demonstrating clearly the utility of using ball milling to obtain nanocrystalline LiH. Furthermore, an inspection to the trend in Fig. 4 shows that the average crystallite size first decreases rapidly, but then slowly as the ball milling time increases. The average crystallite size in Fig. 4 appears to level off at about 25 nm after 180 min of ball milling, suggesting that there will be little crystallite size refinement with a more prolonged ball milling. From the shape of the 1 1 1 XRD peaks, the crystallite size distribution can be calculated using Eqs. (6)–(8). Fig. 5 compares the crystallite size distributions of LiH powders

Fig. 4. The surface-weighted average crystallite size of LiH as a function of ball milling time (close dots) and thermal exposure (open dot). The solid and dashed lines are guides for the eyes.

with different ball milling times. It can be seen that as the ball milling time increases, the crystallite size distribution becomes sharper and shifts towards smaller crystallite sizes. In particular, the median and the deviation of the lognormal crystallite size distribution are calculated to be 23 nm and 0.65 for the 45 min ball-milled LiH, 14 nm and 0.62 for the 90 min ball-milled LiH and, finally, 12 nm and 0.60 for the 180 min ball-milled LiH. Two additional features are noted from this figure, that is: (i) there are almost no crystallites with sizes lower than 3 nm and (ii) the crystallite sizes extend up to 150, 100, and 75 nm for the 45, 90, and 180 min ball-milled LiH, respectively. It can be thus concluded that the nano-crystallization process of LiH during ball milling proceeds via continuous reduction of the crystallites with sizes larger than about 10 nm, which appears to be the final crystallite size that can be achieved by the specific ball milling condition applied in this study. The internal stresses calculated with the aid of Eq. (5) are zero in all cases, and are therefore not plotted. The absence of internal stresses indicates that the interior of the LiH nano-crystals is free of lattice defects, and therefore free of dislocations. This fact suggests that the ball-milled LiH crystals are either predominantly single crystals with few dislocations or polycrystals with few dislocations in the crystal interior and at grain boundaries. If the ball-milled LiH powder is composed of single crystal particles, it can then be inferred that the crystallite size refinement of the LiH powder (or hereafter called the crystal size refinement since the particles are assumed to be single crystals) in the ball milling process does not occur through severe plastic deformation with dislocation-mediated activities. Instead, the crystal size refinement is achieved via repeated brittle fracture, and the final crystal size is limited by particle sliding past each other during ball collision. This hypothesis can also explain two other phenomena observed in this study, that is: (i) the crystallite size distribution of the LiH powder becomes narrower as the ball milling time increases (Fig. 5) and (ii) the crystallite size decreases very slowly after long milling time (Fig. 4). Highenergy ball milling is regarded as mini-forging of many powder particles trapped between two colliding balls [49]. The collision force is mainly carried by large particles, whereas smaller particles can slide past each other. As a result, large particles fracture and become smaller, while small particles slide past

Fig. 5. The normalized crystallite size distribution of LiH for the different ball milling times, with and without the thermal exposure.

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each other with minimal decrease in size. As such, the distribution of the particle size (i.e., the crystal size because of single crystals) becomes narrower as the ball milling time increases, as shown in Fig. 5. In addition, fracture of small particles (<50 nm) is very difficult because the collision force is counterbalanced by the friction force among nano-particles trapped between the two colliding balls, and the collision energy is consumed by friction energy via sliding of nano-particles with respect to each other. As a result, the crystal size decreases very slowly after long milling time, as shown in Fig. 4. If the ball-milled LiH powder is composed of polycrystals with nano-grains and no lattice bending and dislocations, then the short thermal exposure during ball collision must have effectively imposed an annealing treatment to the powder particles trapped between two colliding balls. The internal stresses and dislocations, if any, are presumably eliminated by this local and short annealing treatment. However, the possibility of this mechanism is very low, considering the simultaneous presence of three facts: (i) zero internal stresses in LiH crystals in all cases, (ii) the non-uniform temperature at the collision site during ball milling, and (iii) retention of nano-grains as small as 10 nm even though all dislocations and lattice bending in LiH crystals have been annealed out. Nevertheless, powder morphology examination via SEM and TEM would be extremely useful to provide additional physical evidence to distinguish the mechanisms for the crystallite size refinement of LiH powder. This will be discussed in Sections 3.3 and 3.4.

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LiH are not deformed and the effect of the thermal exposure would be, if any, to relax the deformation. With respect to the peak sharpening observed after the thermal exposure, this cannot be caused by release of internal stresses because the LiH crystallites generated by ball milling are free of lattice defects, and thus it is concluded that crystal growth has taken place during the thermal exposure. This conclusion can be confirmed by examining Figs. 4 and 5, which show, respectively, that the average crystallite size increases by a factor of two and that the crystallite size distribution shifts towards higher crystallite sizes after the thermal exposure. However, note that the crystallites in the thermally exposed LiH are still nanometer-sized, clearly indicating the ball-milled LiH is relatively stable even after exposure to a temperature as high as 0.58Tm (where Tm = 680 ◦ C is the melting temperature of LiH). 3.3. SEM and SSA analyses Fig. 7 shows the SEM images of LiH powder ball-milled for 180 min. Compared with the morphology of the as-purchased

3.2. XRD analysis of the thermal exposure effect Shown in Fig. 6 are the 1 1 1 XRD peaks of the 180 min ballmilled LiH with and without thermal exposure at 285 ◦ C. Two distinct features are noted in this figure, one being the absence of peak shifting and the other the existence of peak sharpening after the thermal exposure. The similitude in the position of the two peaks indicates that the lattice parameter is the same (see Fig. 3), as one would expect because the crystallites in the ball-milled

Fig. 6. 1 1 1 XRD peaks for the 180 min ball-milled LiH with (dotted line) and without (solid line) thermal exposure after correction by instrumental shifts. For comparison, the two peaks have been normalized by imposing the same maximum intensity.

Fig. 7. SEM images of LiH powder with 180 min of ball milling at two different magnifications with: (a) showing the overall view of the agglomerated powder and (b) a close view of two or three agglomerates and the presence of some nano-particles forming these agglomerates.

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Fig. 8. (a) TEM bright-field image of the ball-milled LiH powder (the agglomerate of Type I), showing agglomeration and some degrees of cold welding of individual particles and (b) the corresponding dark-field image, indicating that the individual crystal size ranges from 6 to 30 nm.

powder (Fig. 1), powder particles have become much finer for the 180 min ball-milled powder. Agglomeration of fine particles is also obvious for the ball-milled powder. Some of the individual nano-particles forming the agglomerates appear to be visible in Fig. 7(b). However, given the resolution limit of SEM, it is hard to conclude that these nano-particles are made of one crystal or several smaller nano-particles. Furthermore, some particles cannot be distinguished between hard agglomerates and polycrystalline particles. Detailed TEM analysis would be necessary to shed light on these issues. The specific surface area of the as-purchased LiH powder is found to be 4.61 m2 /g, whereas the corresponding value for the 180 min ball-milled powder is 13.10 m2 /g. If LiH particles are assumed to be spherical, the equivalent particle size can be calculated from the SSA. The equivalent particle diameter so calculated is 1.58 ␮m for the as-purchased LiH powder and 0.56 ␮m for the 180 min ball-milled LiH powder. Note that the equivalent particle diameter of the as-purchased LiH is in excellent agreement with the SEM examination, indicating that the as-purchased LiH particles are loosely agglomerated as shown in Fig. 1. However, the equivalent particle diameter of the 180 min ball-milled powder is substantially different from the XRD analysis which shows the crystallite size in the order of 20 nm. The discrepancy suggests that the ball-milled LiH powder is either polycrystals with nano-grains or nanometer-sized single crystals with strong agglomeration. The latter case is supported by the SEM observation (Fig. 7); however, SEM observations do not rule out the possibility of the presence of polycrystals with nano-grains. TEM analysis would be necessary to clarify this.

relatively smooth surfaces. The crystal sizes of the primary particles forming the agglomerates range from 6 to 30 nm, as revealed from the dark-field image (Fig. 8(b)). Another type of the morphology (called Type II hereafter) is shown in Fig. 9. The agglomerates of this type contain fewer pores and have faceted surfaces. The size of the facet ranges from 10 nm to more than 50 nm, and is believed to correspond to the sizes of the primary particles. The agglomerates of Type I are likely formed by a loose powder agglomerate which experiences a short thermal excursion during ball milling. As pointed out previously, the local temperature at the collision site at the moment of colli-

3.4. TEM analysis TEM analysis indicates that the ball-milled LiH powder is highly agglomerated, exhibiting two types of morphologies. Fig. 8 shows one of the agglomerate morphologies (called Type I hereafter). These agglomerates contain many pores, but have

Fig. 9. TEM bright-field image of the ball-milled LiH powder (the agglomerate of Type II), showing agglomeration, some degrees of cold welding, and faceted surfaces of individual particles. The size of the facet ranges from 10 nm to more than 50 nm.

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sion can be as high as 500 ◦ C during ball milling [32,33]. Such a thermal exposure, however, is very short because the loading and unloading time during one collision is in the order of 10−5 s [47]. The smooth surfaces observed can result from such a short thermal exposure, and the retention of nano-grains is likely due to the presence of pores between primary particles and the short duration of thermal excursion. The agglomerates of Type II, in contrast, are likely formed through short thermal excursion in the presence of pressure, followed by subsequent fracture during ball milling. The presence of pressure during ball milling results in denser agglomerates, whereas pores between primary particles prevent crystal growth during the short thermal exposure. The conclusions drawn from the TEM analysis of the ballmilled LiH powder are in good agreement with XRD, SEM, and SSA analyses discussed in Sections 3.1 and 3.3. First, the TEM analysis confirms that the primary particle sizes range mainly from 6 to 50 nm, as quantified via the XRD analysis (Figs. 4 and 5). The TEM analysis also unambiguously reveals strong agglomeration between primary particles. Because of such strong agglomeration, the SSA of the ball-milled LiH powder would be substantially lower than that predicted based on the summary of the surfaces of primary particles, and thus the TEM analysis is consistent with the SSA measurement. The TEM analysis also supports the repeated brittle fracture as the mechanism for the crystal size refinement of LiH powder because the ball-milled LiH powder is not composed of solid particles with nano-grains. Solid particles with nano-structures are the required evidence for the crystallite size refinement via severe plastic deformation with dislocation-mediated activities, as found in many ball-milled metals [35,36,48]. Finally, the TEM analysis is also consistent with the conclusion from the XRD analysis that the ball-milled LiH crystals have no internal stresses and no dislocations. Since the crystal size refinement is achieved via repeated brittle fracture, there is no need for dislocation activities. Furthermore, dislocations, if any, would escape to the

303

surfaces of nano-crystals readily because of the image force on the other size of the free surface [49]. Since the strong LiH agglomerates are formed from nano-particles via short thermal exposure during ball collision, there will be no dislocations and lattice bending at the particle boundaries. Taken together, XRD, SEM, TEM, and SSA analyses reveal that the ball-milled LiH powder is composed of strong agglomerates which are made of nano-crystals free of lattice defects and bending. It is noted that both types of the morphologies observed in the as-milled condition have been inherited by the thermally exposed samples. Fig. 10 shows the agglomerate of Type I after the thermal exposure at 285 ◦ C for 1 h. In comparison with the same type of agglomerates before the thermal exposure, two features are noted. The first is the smoother surface of the thermally exposed agglomerates. The second is the slightly larger crystal sizes for the thermally exposed agglomerates, indicating that the crystal growth has taken place during the thermal exposure. In fact, in some regions of the agglomerate many small crystals are in the process of linking together to form a large crystal, as indicated by the ellipse in Fig. 10(b). Such crystal growth phenomenon is even more obvious in the agglomerate of Type II. As shown in Fig. 11, many small particles within an agglomerate of ∼0.8 ␮m have grown together to form a large crystal. Smoother surfaces are also obvious for the thermally exposed powder when compared with the same type of agglomerates before the thermal exposure. Finally, it is noted that in spite of smoother surfaces and growth of crystals, pores are still present in the thermally exposed sample. Based on the TEM analysis, it can be concluded that grain growth has taken place during the thermal exposure at 285 ◦ C. However, most of the crystals remain in nano-scales. All of these are in good accordance with the XRD analysis (Figs. 4 and 5). The retention of nano-crystals and sintering of agglomerates far from completion indicate that the ball-milled LiH powder is relatively stable even after exposure to a temperature as high as

Fig. 10. (a) TEM bright-field image of the thermally exposed LiH powder (Type I), showing some smoothening of the surface of the agglomerate in comparison with the surface of the ball-milled powder and (b) the corresponding dark-field image, indicating that most of the individual crystals have sizes ranging from 6 to 35 nm. The circled region appears in the process of growing into a large crystal with a size of about 100 nm.

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Fig. 11. (a) TEM bright-field image of the thermally exposed LiH powder (Type II) and (b) the corresponding dark-field image, showing that many individual crystals are in the process of linking together to form a large crystal, as indicated by the ellipse.

0.58Tm for 1 h. Such a good thermal stability is presumably due to the presence of many pores between primary particles at the ball-milled condition, and expected to provide good stability in the hydrogen sorption and desorption cycles when mixed with LiNH2 . 4. Concluding remarks The present set of study investigates the powder characteristics of LiH crystals as a function of high-energy ball milling and subsequent thermal exposure at 285 ◦ C. The following conclusions can be offered based on this study: 1. High-energy ball milling at room temperature can effectively reduce LiH crystals to an average size of 25 nm. 2. The crystal size reduction proceeds rapidly at the early stage of ball milling, whereas the size reduction rate decreases substantially beyond 90 min of ball milling and the average crystal size appears to level off at ∼25 nm after 180 min of ball milling. 3. The crystal size distribution of LiH powder becomes sharper and shifts towards smaller sizes as the ball milling time increases. For the 180-min ball-milled LiH powder, the largest crystals are ∼75 nm with most of crystals having sizes lower than 50 nm and no crystals smaller than 3 nm. The corresponding median and the deviation of the lognormal crystal size distribution are 12 nm and 0.60, respectively. 4. The nano-crystals of ball-milled LiH powder are highly agglomerated. Two major types of agglomerate morphologies are present with one containing many pores and having relatively smooth surfaces (Type I) and the other containing fewer pores and having faceted surfaces (Type II). The crystal sizes of the primary particles forming both types of agglomerates center around 6–50 nm. 5. The crystal size reduction of LiH powder is mainly accomplished via repeated brittle fracture during ball milling. The

formation of many hard LiH agglomerates is assisted by the short thermal exposure during ball collision. Both the brittle fracture and short thermal exposure appear to contribute to the features of no lattice defects and bending in the nano-crystals of ball-milled LiH powders. 6. Both types of agglomerate morphologies in the ball-milled condition are inherited by the thermally exposed samples. Thermal exposure at 285 ◦ C for 1 h results in grain growth and some sintering of agglomerates. However, most of the LiH crystals (>90%) remain in nano-scales (<100 nm) after the thermal exposure. 7. The retention of LiH nano-crystals during the thermal exposure at 285 ◦ C (equivalent to 0.58Tm ) indicates that the ball-milled LiH powder is relatively stable and is ascribed to the presence of many pores which prevent crystal growth in the agglomerates. Acknowledgements This work was supported under the U.S. Department of Energy (DOE) Contract No. DE-FC36-05GO15008. The vision and support of Dr. Carole Read, DOE Technology Manager, is greatly appreciated. The authors are also grateful to Dr. Roger Ristau in assisting the attainment of the TEM pictures. References [1] R.-M. Ren, Z.-G. Yang, L. Shaw, Scripta Mater. 38 (5) (1998) 735–741. [2] R.-M. Ren, Z.-G. Yang, L. Shaw, Nanostruct. Mater. 11 (1) (1999) 25–35. [3] Z.-G. Yang, R.-M. Ren, X.-Q. Xie, L. Shaw, Metall. Mater. Trans. 30A (4) (1999) 1109–1117. [4] Z.-G. Yang, R.-M. Ren, L. Shaw, J. Am. Ceram. Soc. 83 (8) (2000) 1897–1904. [5] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Compd. 217 (1995) 245. [6] F. von Zeppelin, H. Reule, M. Hirscher, J. Alloys Compd. 723 (2002) 330–332. [7] S. Orimo, H. Fujii, K. Ikeda, Acta Mater. 45 (1997) 331.

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