3.1. Microstructural Characterization
The elemental microanalysis measured a C content of 5.07 ± 0.21 vol.% for the samples with 5 vol.% e-GNP nominal content, which means that no losses of ceramic or GNP took place during the powder processing.
The Raman spectra of the processed composite powders (after planetary ball milling) were normalized to the characteristic G band (Figure 1
All the spectra present characteristic D and G peaks centered at ~1350 and ~1580 cm−1
, respectively, as well as the splitting of 2D Raman-active bands centered at ~2700 cm−1
. However, there were remarkable differences between the spectrum of the as-received GNP and the spectra of the composite powders after high-energy planetary ball milling. A significant increase in the D band can be observed in the spectra of the powders. Moreover, other peaks previously described in literature [31
] as related to defects in graphene and carbon-based materials can be detected. These bands, referred to as D* (~1100–1200 cm−1
), D” (~1500–1550 cm−1
), and D’ (~1610–1620 cm−1
) were almost impossible to detect in the spectrum of the as-received GNP, but were clearly visible in the spectra of the powders. Accordingly, all the first-order spectra were fitted to two Gaussian (D* and D”) and three pseudo-Voigt (D, G, and D’) functions.
A significant increase in the intensity (integrated area) ratios of all the defect-related peaks with respect to the G band was observed for the composites powders when compared to the as-received GNP (Table 1
). Nevertheless, the increase was more remarkable for the composite with 1 vol.% e-GNP in all the cases. The drastic increase in the ID
ratio may indicate a reduction in the e-GNP lateral size due to the fracture of the nanoplatelets provoked by the planetary milling, although it is not possible to discard the promotion of structural defects in the GNP during milling. Moreover, the increase in the ID’
ratio would point to an exfoliation process resulting in a lower thickness of the GNP after milling. The D and D’ bands are related in the literature to defects related to graphene layer edges and surface graphene layers, respectively [33
]. A drastic lateral size reduction effect in the GNP due to processing with planetary ball milling was also observed in a previous work for composites with 10 vol.% e-GNP [8
]. The variation in the ID
ratio for the powders with different e-GNP content is also meaningful. The composite powders with 1 vol.% e-GNP exhibited the highest ID
ratio, which may indicate that ball milling is more effective in lateral size reduction of the GNP with a higher ceramic/GBN ratio. Milling of compositions with higher GNP content was less aggressive since the GNPs possess a lubricating effect [7
Regarding the D” and D* peaks, which were previously related to the presence of amorphous carbon and sp2
bonds, respectively [33
], an increase in the intensity ratios was observed for both powders, being more significant for the composite powders with 1 vol.% e-GNP. This confirms that structural damage was introduced in the GNP during high-energy milling.
When analyzing the second-order spectra, a change in the shape of the 2D band was clearly observed on the spectra of the composites powders when compared to that of the as-received GNP. A fitting of the 2D band using two Lorentzian functions could be performed for the spectra of the as-received GNP (Figure A1
a, Appendix A
), which reveals that the nanoplatelets were formed by stacking of more than 10 graphene layers that present a 2D band similar to the one of graphite [34
]. After milling of the powders, the shape of the 2D band was modified and did not correspond to the sum of two peaks. However, the altered shape of the spectra (probably related to the presence of structural defects) did not allow the deconvolution of peaks in order to establish the number of graphene layers in the nanoplatelets after the composite powder processing.
The measured density of the composite with 1 vol.% e-GNP was 5.97 ± 0.05 g/cm3
(99% relative density), while a similar density of 5.81 ± 0.10 g/cm3
(100% relative density) was obtained for the composite with 5 vol.% e-GNP. These values correspond to fully dense composites, and they are analogous to the densities measured in composites densified by the same SPS technique but obtained using milder powder processing techniques [26
The Raman spectra of the sintered ceramic composites are presented in Figure 2
. The intensity ratios obtained after fitting the first-order spectra to the five functions previously described (Table 1
) reveal that, after sintering, the values for the two composites were very similar to each other. For the D* and D” peaks, the intensity ratios were also very close to those obtained for the as-received GNP. This reveals a recovery of the structural damage that was induced during milling of the powders, since it was reported that the ID”
ratio decreases as the crystallinity increases [35
]. It was also found that the restoration of the graphene structure above 1000 °C is related to the smoothing of the ripples and roughness of the flakes, producing a drastic peak width decrease and 2D intensity enhancement [37
]. On the other hand, the ID
ratios were again higher than for the as-received GNP, confirming that the milling of the powders promotes a decrease of the lateral size and an exfoliation of the nanoplatelets. For the composite with 1 vol.% e-GNP, a decrease in ID
until reaching the same value as for the composite with 5 vol.% e-GNP was observed. The former had higher structural damage after milling, which was recovered during sintering of the composite.
The decrease in the number of structural defects also had an effect on the shape of the second-order spectra, which allowed the fitting of this part of the spectra for the sintered composites. The 2D band could be fitted to three Lorentzian functions (Figure A1
b,c, Appendix A
), revealing that the e-GNP filler presented a number of layers lower than 10 [34
] in both sintered composites. Thus, it was confirmed that the high-energy milling provoked an exfoliation of the GNP during the powder processing, turning the nanoplatelets into multi-layered graphene.
The SEM micrographs of the composite polished surfaces annealed in air (Figure 3
) show the presence of small voids dispersed next to the ceramic grains, in a higher number for the composite with the highest e-GNP content. These voids were probably produced by the combustion of the e-GNP during annealing in the presence of oxygen, explaining the greater number of voids in the composite with 5 vol.% e-GNP. However, some of them could also correspond to ceramic grains that were removed by pull-out during the grinding and polishing steps prior to the annealing.
A slight grain refinement can be observed in the composites (0.25 ± 0.11 and 0.20 ± 0.08 μm for 1 and 5 vol.% e-GNP, respectively), compared to the monolithic 3YTZP [26
], in agreement with the ceramic grain growth inhibition effect of GNP reported in the literature [26
]. The ceramic grains were in both cases almost equiaxed.
The low-magnification SEM micrographs of the polished cross-sections of the composites acquired using BSE (Figure 4
) show a continuous light-gray area, corresponding to the 3YTZP matrix, and a discontinuous dark area, corresponding to the e-GNP. The e-GNP distribution achieved was quite homogeneous, and the planetary milling effect was very strong on the GNP size (Figure 4
). Only some GNP aggregates with micrometric lateral size can be seen (Figure 4
a), while the rest were broken and reduced to submicrometric size. In the composites with higher e-GNP content (5 vol.%; Figure 4
b), we can observe a higher number of platelet-like e-GNPs, in agreement with the Raman results, which points to a less aggressive milling due to the lubricating effect of the GNPs. In this composite, the e-GNPs surround the ceramic grains creating a contouring effect. These e-GNPs, therefore, adopt the random orientation of the grain boundaries, as reported for higher e-GNP content [8
], in contrast with the GNP preferential orientation (perpendicular to the pressing axis of the SPS) typically reported for ceramic composites prepared from powders not subjected to high-energy milling [26
]. These observations indicate a microstructural isotropy of the composites from powders processed by planetary ball milling, with the e-GNPs surrounding the ceramic grains.
3.2. Mechanical Characterization
The results of the flexural strength of the tested materials are displayed in Table 2
. The maximum flexural strength value was measured in a monolithic 3YTZP test specimen. However, considering the average performance, the composite with 1 vol.% e-GNP showed slightly superior results, exhibiting a 7% flexural strength increase with respect to the monolithic ceramic. The scarce agglomerates observed did not seem to have a negative effect on the flexural strength of the material reinforced with 1 vol.% GNP. The flexural strength decreased abruptly (~50% decrease) when the e-GNP content increased from 1 to 5 vol.%. Another remarkable fact is that the scatter of the data was lower in the composites compared to the monolithic ceramic, indicating a more reliable behavior in the composites. This is in agreement with the uncertainties observed in the data of flexural strength of similar composites by other authors [9
], and it can be interpreted as a decrease in the brittleness of the 3YTZP ceramic as a consequence of the e-GNP addition. Our results of flexural strength for the monolithic 3YTZP are lower than previously reported ones [9
]. For the composites, the absolute values of flexural strength are slightly higher than those reported by Obradovic and Kern [9
] and coincident with those reported by Li et al. [22
]. A decrease in flexural strength upon increasing the GNP content similar to our results was reported for 3YTZP [9
], as well as for a Si3
]. However, for this last matrix and rGO filler, a different trend was reported, with an increase in flexural strength with filler content up to 4 vol.% rGO [11
The flexural modulus of the 3YTZP ceramic matrix, estimated from the slope of the first portion of the flexural stress–strain curves as described in ASTM D-790, increased by about 10% with the addition of 1 vol.% e-GNP (see Table 2
). In contrast, the addition of 5 vol.% e-GNP resulted in a decrease in the flexural modulus. A remarkable fact is that the experimental uncertainty was much lower (an order of magnitude) in the composite with 1 vol.% e-GNP compared to the monolithic 3YTZP and to the 5 vol.% e-GNP composite. The low uncertainty value indicates a high reproducibility of the flexural modulus results and, therefore, an increased reliability of the composite with 1 vol.% e-GNP in bending conditions, in agreement with the results of flexural strength.
Although the elastic modulus under uniaxial stress conditions is equivalent to the flexural elastic modulus, our results show systematically higher values for the Young’s modulus measured by the acoustic technique (Figure 5
). The behavior of the flexure modulus also differed from the Young’s modulus for low e-GNP contents. The Young’s modulus, E, decreased continuously with increasing e-GNP content (Table 3
). The decrease in the elastic modulus with increasing GBN content was already reported for 3YTZP composites with GBN in our previous studies [39
], and it can be attributed to the effect of adding a more elastic phase (GBN filler) to a rigid matrix (3YTZP ceramic in this case).
The hardness in the composites decreased with increasing e-GNP content, which is a well-known effect in these ceramic composites and a direct consequence of introducing a soft phase into a harder ceramic matrix. No mechanical anisotropy was detected for the composite with 5 vol.% e-GNP. The hardness values in this composite were identical to those reported for composites prepared without high-energy milling [26
], but the composite with 1 vol.% e-GNP exhibited a slightly lower average hardness than the correspondent composite processed with ultrasonic mixing [26
]. This decrease in hardness for the composite with a low content of exfoliated graphene nanoplatelets could be due to the more aggressive milling in this composite, as previously described in the Raman analysis.
The Vickers indentation imprints were well defined on the top surface of the studied composites (Figure 6
), although some isolated cases of spalling were observed in composites with 5 vol.% e-GNP (not shown). The cracks arising from the corners of the imprints were quite straight, due to the small size of the e-GNP, which minimized deflection or deviation of the cracks. The crack length diminished in composites with increasing e-GNP content (see Table 4
), and crack tortuosity increased slightly (Figure 6
b,d), suggesting that a certain small crack propagation inhibition effect was exerted by the e-GNP.
The same crack behavior (shape and length) was observed when indentation was applied on the cross-sections of the composites (see Figure 6
and Table 4
). This indicates an isotropic response to the crack propagation of these composites. The effect of increasing the e-GNP content was again a slight decrease in the crack length. This decrease was mainly caused by toughness enhancement mechanisms such as crack deflection by the e-GNP (Figure 7
), which relieved the stress on the crack tip, thereby reducing the effective crack length. This mechanism was more effective with higher e-GNP content; thus, it was responsible for the increased tortuosity of the cracks observed in the composite with 5 vol.% e-GNP (Figure 7
). The fracture toughness estimated by DCM using the Anstis model did not follow a clear trend with the e-GNP content. The Anstis equation was chosen since it was the most restrictive model, giving the lowest values for fracture toughness. Niihara’s equation gave higher values, showing a similar evolution of the fracture toughness with the GNP content. Shetty’s equation gave values between those for Anstis and Niihara with similar fracture toughness for both composites, slightly lower than the monolithic 3YTZP. Although these estimations of fracture toughness by DCM should not be taken as accurate values, they can give us a rough approach.
Additional cracks around the indentation imprints, just outside the plastic deformation zone, could be observed at high magnification in all the indented specimens (insets in Figure 6
a,b). They correspond to shallow lateral cracks [40
], a variant secondary crack system which has been reported in ceramic coatings on soft substrates. While lateral cracks in homogeneous materials develop typically well below the surface during unloading at high indentation loads, in non-homogeneous specimens with a surface hardened layer, the lateral crack may initiate much closer to the surface [41
]. We may consider our composites as arrays of hard ceramic layers on top of softer GNP substrates. Since lateral cracks can cause chipping (when they nucleate and divert toward the surface), a shallower lateral crack formation would reduce the material removal rate during wear, increasing wear resistance [41
]. Future studies will be carried out to check the tribological response and the fracture toughness of these composites.