The average crystallite size of Si particles was calculated by Scherrer formula from XRD patterns, and the results of Si
L, Si
S, Si
L+M, Si
S+M and Si
S+MM are listed in
Table 3 [
34]. The Raman spectra of the 5 samples were analyzed by decomposing the spectra into the crystalline and amorphous bands. The Raman bands at near 510 cm
−1 and 480 cm
−1 were attributed to phonon modes of a-Si and c-Si, respectively, and the ratio of their integrated intensities could be used to estimate the amorphization degree because other bands at near 320, 430, 600 cm
−1 were the LO, TA and combination modes, respectively, which was two orders smaller in magnitude than TO mode in the crystalline phase. Gaussian fitting was performed on the Raman peaks, near 480 cm
−1 and near 510 cm
−1. By calculating the integral intensity of curves, the amorphization degree estimated by Raman (
) was obtained from Equation (1) [
35,
36]:
where
and
are the integral intensity of a-Si and c-Si, respectively.
is the ratio of Raman scattering cross-section of crystalline and amorphous components, and here, a typical value of
is used. The analysis results of the 5 samples are shown in
Table 3.
As a comparison, XRD pattern fitting was used to estimate amorphous ratio [
37,
38], and the curves of a-Si and c-Si at 28.4°, 47.3° and 56.1° were obtained, respectively. By calculating the integral intensity of curves, the amorphization degree estimated by XRD (
) was obtained from Equation (2):
where
and
are the integral intensity of a-Si and c-Si in XRD, respectively.
is a constant coefficient, and
is usually used for elementary substances. The analysis results of the 5 samples are shown in
Table 3. New Si particles reacted with H
2O at high temperatures during spark erosion to form a SiO
2 layer on the surface, and the fresh Si particles’ surface generated by collision during bead milling reacted with alcohol, forming oxides [
39,
40]. The oxygen contents of the 5 samples are presented in
Table 3.
3.2.1. Internal Structure of Si Micro/Nano Particles Prepared by Spark Erosion
The X-ray diffraction patterns and the Raman spectra of Si
L and Si
S are shown in
Figure 8a–b, respectively. Diffraction peaks of Si
L and Si
S can be well indexed to the cubic Si phase (JCPDS. Card no. 01-0787) with (111) at 28.4°, (220) at 47.3°, (311) at 56.1°, (400) at 69.1° and (331) at 76.4°, as shown in
Figure 8a. The intensity of the diffraction peak of Si
L was higher than that of Si
S, and the peak width was narrower, indicating that Si
L had a higher crystallinity and a larger crystallite size, which was in agreement with the resulted from Raman spectra. The average crystallite sizes of Si
L and Si
S were calculated by Scherrer formula from XRD patterns to be 68.92 nm and 26.47 nm (
Table 3).
As shown in
Figure 8b, the peak at near 510 cm
−1 corresponds to crystalline Si (c-Si), and the peak intensity of Si
L was obviously stronger than that of Si
S, indicating that the degree of crystallization of Si
L was higher than Si
S. A weak peak at near 480 cm
−1 for both samples was assigned to amorphous Si (a-Si), indicating that the amorphous phase had already existed inside both samples. As shown in
Figure 8c, Red curves were fitting curves of a-Si, and blue curves were those of c-Si. Curves in other colors in
Figure 8c are peaks at other Raman shifts and the intensity, which was not used to calculate the amorphization degrees in Equation (1). By fitting the Gaussian peak of the Raman curve, combined with Equation (1), the amorphization degrees of Si
L and Si
S were 31.8% and 43.6%, respectively (
Table 3). As shown in
Figure 8d, the amorphization degrees of Si
L and Si
S were also estimated by fitting XRD patterns with Equation (2), and the results were 27.8% and 46.3% (
Table 3), which agree well with those calculated by Raman.
Figure 9 shows HRTEM images of Si
L and Si
S. The nanocrystals are marked with a white circle, and the orientations are labeled by white lines. It can be verified from
Figure 9b,d that the crystallite size of Si
S was smaller than that of Si
L and the crystal orientations were disordered in Si
S.
The structure of SiO
2 was amorphous, and the oxygen contents of Si
L (1.3%) and Si
S (2.5%) were both low (
Table 3), which means that the dense SiO
2 layer prevented the inner Si from being oxidized further. The oxygen content of Si
S was higher than that of Si
L because Si
S had a smaller particle size and, therefore, had a larger specific area than Si
L. It could be deduced from these results that small pulse duration not only manufactures smaller particles but also created the smaller crystals embedded in a highly amorphous matrix.
During spark erosion, after vaporizing or melting, the silicon particles are either in a gaseous or a liquid state. Their internal structure changes from a single crystal structure to an amorphous structure, in which the silicon atoms are arranged irregularly. The amorphous structure contains higher potential energy, as the distance between the Si atoms is not equal to the distance in the equilibrium. During the condensation process, the amorphous structure recrystallizes in a short time and transforms into the crystalline structure that contains the smallest internal energy and is the most stable. The cooling rate has a very important influence on the size of crystallites [
41,
42]. Remaining at a high temperature for a long time will slow down the cooling rate of the formed particles, which will increase the time of the recrystallization process, resulting in forming a small number of crystal grains inside the particles with a larger size. As a result, the short pulse duration leads to short high-temperature time, which increases the cooling rate and the nucleation rate of the formed particles during recrystallization, as well as the number of crystal grains that have smaller size and higher disordered crystal orientations [
43].
3.2.2. Internal Structure of Si Nanoparticles Prepared by Bead Milling
The X-ray diffraction patterns and the Raman spectrum of Si
L+M, Si
S+M and Si
S+MM are shown in
Figure 10. The average crystallite size of Si
L+M, Si
S+M and Si
S+MM is calculated by the Scherer formula from XRD patterns being ~11.73 nm, 6.18 nm and 4.23 nm, respectively (
Table 3). The peak intensity of Si
L+M was higher than that of Si
S+M, but the peak width of Si
S+M was wider, indicating that the smaller the crystallite size of the particles before bead milling was, the smaller it was after bead milling. Comparing Si
S+M and Si
S+MM in
Figure 10a, with the increase of the milling time, the intensity of the diffraction peak of Si
S+MM was further weakened, and the peak width was further broadened, indicating that increasing the milling time could further refine the grains and increase the amount of amorphous phase.
This conclusion is further verified in the Raman spectrum. As shown in
Figure 10b, the c-Si peak of Si
L+M near 510 cm
−1 was stronger than that Si
S+M. Increasing the milling time, the c-Si peak of Si
S+MM becomes weaker and weaker. Simultaneously, the peak of a-Si near 480 cm
−1 increased for all the samples with Si
S+MM > Si
S+M > Si
L+M. According to Equation (1) and
Figure 10c, the amorphization degrees (
) of Si
L+M, Si
S+M and Si
S+MM were calculated to be 69.4%, 74.8% and 83.7%, respectively (
Table 3). Similarly, the amorphization degrees (
) of these three samples estimated by Equation (2) and
Figure 10d were 70.4%, 75.6% and 83.1% (
Table 3). Through the analytical results of all 5 samples, we could conclude that the amorphization degrees estimated by fitting Raman spectra were highly similar to those calculated by fitting XRD patterns, which means the results are highly credible.
Figure 11 shows HRTEM images of Si
L+M, Si
S+M and Si
S+MM. The nanocrystalline structure embedded in the amorphous structure was formed inside. As a comparison to the raw material of high-energy ball milling and two-step bead milling (0.8, 0.1 mm ZrO2 beads), the internal structure of Si micro/nanoparticle was different and easier to form amorphous regions because these silicon particles were either in a gaseous or a liquid state during spark erosion [
27,
28]. For the same grinding time, Si
S+M had a smaller grain size and a higher degree of amorphization than Si
L+M, which was mainly due to the difference in raw materials (Si
L and Si
S); however, the difference between Si
L and Si
S had become smaller, indicating that bead milling for a long time could weaken the influence by raw materials. For the same raw material (Si
S), Si
S+MM with longer bead milling time had a smaller grain size and a higher degree of amorphization than Si
S+M, indicating that prolonging the grinding time could not significantly reduce the size of silicon particles, but it could further refine the crystal grains and form more amorphous structures.
However, the increase in grinding time also increased the degree of oxidation of silicon nanoparticles. The degree of oxidation evolved depending on the specific surface area of the particles and the bead milling time. The bead milling time of Si
L+M and Si
S+M was the same, but the increased oxygen content of Si
S+M (6.2%) was slightly higher than Si
L+M (5.8%) because of the smaller size of Si
S+M. As the grinding time of Si
S+MM was 4 h longer than Si
S+M, the oxygen content of Si
S+MM (8.7%) was obviously higher than Si
S+M (6.2%) (
Table 3). Due to the repeated collisions of beads, when the receiving energy of a silicon atom in the grain of a silicon particle was greater than its bond energy, the distance between the atoms increased or decreased, and the potential energy increased. The atomic bond was broken to generate holes or lattice shifts, and the original parallel hexahedral structure was changed. The big grains were split into multiple smaller grains and even formed an amorphous structure with higher potential energy [
33]. The greater the number of collisions, the smaller the crystallite size and the higher the proportion of the amorphous structure [
44]. These effects were gradually transferred from the surface to the interior of the particles and gradually form a structure of Si nanocrystals embedded in amorphous Si.
This technical route combined the advantages of spark erosion and bead milling, e.g., spark erosion saved much time and energy for bead milling. From the observation of the morphology, internal structure and amorphization degree of Si particles, we could control the formation of nanocrystal and amorphous regions. Both spark erosion and bead milling produced defects inside the silicon particles, and bead milling could produce more defects than spark erosion because the bead milling takes a longer time and does not have a recrystallization process. Repeated collisions during the bead milling process resulted in the formation of a large number of point defects (vacancies, etc.), line defects (mixed dislocations), surface defects (grain boundaries, twin boundaries, etc.) in the particles, and the resultant Si particles were amorphous in a high degree. The processing parameters of spark erosion and bead milling had an important impact on the structure of resultant Si nanoparticles. The yield of this route to get Si nanoparticles depends more on spark erosion than on bead milling, as the yield of bead milling could be easily increased through the optimization of production equipment. The micro/nano Si particles could be collected as waste materials from the factory where Si bulk materials are processed by spark erosion, such as punching, slotting, etc.