F.
Picollo
ab,
A.
Battiato
b,
F.
Bosia
ac,
F.
Scaffidi Muta
a,
P.
Olivero
*ab,
V.
Rigato
d and
S.
Rubanov
e
aPhysics Department and “NIS Inter-departmental Centre”, University of Torino, Torino 10125, Italy. E-mail: paolo.olivero@unito.it
bNational Institute of Nuclear Physics, Section of Torino, Torino 10125, Italy
cApplied Science and Technology Department, Politecnico di Torino, Torino 10129, Italy
dNational Institute of Nuclear Physics, National Laboratories of Legnaro, Legnaro 35020, Italy
eIan Holmes Imaging Centre, Bio21 Institute, University of Melbourne, Victoria 3010, Australia
First published on 9th June 2021
Carbon exhibits a remarkable range of structural forms, due to the availability of sp3, sp2 and sp1 chemical bonds. Contrarily to other group IV elements such as silicon and germanium, the formation of an amorphous phase based exclusively on sp3 bonds is extremely challenging due to the strongly favored formation of graphitic-like structures at room temperature and pressure. As such, the formation of a fully sp3-bonded carbon phase requires an extremely careful (and largely unexplored) definition of the pressure and temperature across the phase diagram. Here, we report on the possibility of creating full-sp3 amorphous nanostructures within the bulk crystal of diamond with room-temperature ion-beam irradiation, followed by an annealing process that does not involve the application of any external mechanical pressure. As confirmed by numerical simulations, the (previously unreported) radiation-damage-induced formation of an amorphous sp2-free phase in diamond is determined by the buildup of extremely high internal stresses from the surrounding lattice, which (in the case of nanometer-scale regions) fully prevent the graphitization process. Besides the relevance of understanding the formation of exotic carbon phases, the use of focused/collimated ion beams discloses appealing perspectives for the direct fabrication of such nanostructures in complex three-dimensional geometries.
Moving towards more “defective” and technologically viable forms of sp3-bonded carbon, different forms of polycrystalline diamond,9 ultra-nanocrystalline diamond,10 nano-twinned diamond11 and amorphous diamond-like carbon12,13 have been widely investigated for several decades, with the promise of further expanding the applicability of extreme physical properties into technological landscapes in which synthesis and fabrication techniques can be realistically scaled to large production volumes. In this context, the higher thermodynamical stability of sp2-bonded carbon at room pressure and temperature conditions represents a fundamental limitation: in these conditions, graphite and graphite-like phases constitute the ultimate “ground state” for carbon structures when a critical amount of structural disorder is introduced. For this reason, substantial efforts have been made in the synthesis of amorphous carbon phases characterized by a high fraction of sp3 bonds,14–16 but the pursuit of a 100% fully sp3-bond amorphous carbon phase is still ongoing. A careful control of environmental parameters (pressure in particular) allows the engineering of novel forms of carbon, as demonstrated by the fact that exerting high (i.e. ∼102 GPa) pressures on glassy carbon (i.e. an amorphous sp2 phase) yields the formation of phases characterized by high sp3 content with no long-range ordering, whose structural stability can to some extent be tuned if an equally careful control of temperature variable can be achieved.17–19 In this context, a powerful and versatile tool is represented by the local laser heating of different types of carbon structures under different mechanical stress conditions, either exerted from external pressure sources20 or established within the sample by the coexistence of carbon phases characterized by different densities and mechanical properties.21
Local laser heating was combined with the possibility offered by MeV ion irradiation to create sub-superficial graphitic structures within bulk diamond thanks to the strongly non-linear damage profile of energetic ions in matter. The ion-damage-induced collapse into a graphitic phase of layers with sub-μm thickness localized within the diamond crystal determines substantial local variations in both atomic density and mechanical parameters (Young's and shear moduli), that can in turn develop strong (i.e. ∼10 GPa) and highly localized internal stresses, without the need of using external pressure sources.22 In these conditions, optical absorption of the laser light at different power densities from the sub-superficial compressed graphitic layers allowed a fine control of local temperature variations, and thus an accurate exploration of the graphite–diamond–liquid triple point.23 In this context, the employment of other types of radiation (e.g. X-ray nano-beams) could be successfully employed to engineer structural damage with high spatial resolution, as already successfully demonstrated in other types of substrates.24–26
More recently, a careful control of the in situ laser-induced heating of glassy carbon kept at high (i.e. ∼50 GPa) pressure by means of a diamond anvil cell allowed the exploration of a very specific (and up to then scarcely studied) portion of the phase diagram of carbon, which resulted in the first demonstration of the synthesis of quenchable fully-sp3 bonded amorphous carbon phase. This stable amorphous phase of carbon was unequivocally demonstrated to be based on a sp2-free structure by means of high-resolution transmission electron microscopy (HRTEM) and electron energy loss spectroscopy (EELS), and exhibited properties of optical transparency, high density and extreme stiffness that were comparable to those of diamond.20
In the present work, we take advantage of a high-resolution lithographic technique based on the use of masked MeV ions to define sub-superficial amorphous nanostructures in the diamond bulk induced by atomic collisions. We demonstrate by means of HRTEM and EELS that these structures are lacking any measurable fraction of sp2 bonds, specifically because their size (i.e. ∼100–200 nm, depending upon fabrication parameters) and depth below the crystal surface (i.e. ∼1.6 μm) is such to inhibit any form of graphitization by the development of strong (i.e. >40 GPa) internal pressures. These results demonstrate for the first time the possibility of direct MeV-ion-beam writing with high spatial resolution a quenchable amorphous phase in diamond at room conditions, with no need of externally applied pressures.
Fig. 1 Sample geometry and profile of the ion-induced structural damage. (a) Schematic representation of the sample geometry: the metallic mask with the nanometric aperture determines the formation of both the continuous shallow layer and of the nanochannel upon MeV ion irradiation. (b) Three-dimensional plot of the cross-sectional profile of ion-induced damage density as resulting from SRIM simulation. (c) Corresponding two-dimensional plot: the size and shape of the region damaged beyond the estimated critical threshold (in red) corresponds to the features observed in Fig. 4a; note that the same plot is reported as an inset of Fig. 4a for sake of comparison with experimental data. |
After MeV ion irradiation and subsequent mask removal, the samples were thermally annealed in vacuum at 950 °C, with the scope of allowing the structural reorganization of the highly damaged buried nano-regions, while removing residual damage from the regions irradiated at intermediate depths.
The nanochannels of highly-damaged carbon phase are expected to form where the structural damage (here parameterized as a volume density of created vacancies, as predicted by the SRIM Monte Carlo simulation code27 in a linear approximation) exceeds a critical threshold, whose value has been estimated as ∼(6.4 ± 1.5) × 1022 cm−3 on the basis of the measured dimensions of the nanostructures (see Fig. 1b, c and 4a).
This value is in good agreement with previous estimations of the parameter, commonly referred to as “graphitization threshold”, in the (5 − 7) × 1022 cm−3 range.28–32 As shown in Fig. 1 (as well as in the inset of Fig. 4a), the SRIM-based model of the damage profile (which also suitably describes the trajectories of laterally straggled ions) accurately predicts not only the ∼1.6 μm depth of the nanochannels below the surface, but also their overall shape.
As shown in Fig. 2a and b, the bright-field TEM cross-sectional micrograph and related selected area diffraction pattern indicate that the as-irradiated microstructures consist of a fully amorphized phase. Remarkably, the same is observed also after the annealing step (see Fig. 2c and d), thus indicating that the thermal process stabilizes the structures without inducing any re-crystallization of either sp2 or sp3 phases. Dark contours are also visible around the nanostructures, indicating lattice strains due to a high local concentration of point defects. The diffraction patterns in Fig. 2b and d only show broad rings that are typical of amorphous structures. It is worth remarking that the radius of the first ring correlates with the positions of the {111} diffraction spots generated from the surrounding crystalline diamond matrix, as reported in the reference diffraction pattern reported in the inset of Fig. 2. This confirms that the probed phase is fully amorphized. The absence in the diffraction patterns of features related to sp2 bonding (i.e. rings corresponding to the {002} lattice plane of graphite) can be attributed to a low fraction of sp2 bonds or to a predominant orientation of graphite basal plains normal to the electron beam direction.
In order to provide direct insight into the nature of the chemical bonds within the amorphized nano-regions, EELS analysis was carried both before and after thermal annealing, in the energy ranges corresponding to the K absorption edge of carbon and the plasmonic energy loss. As far as the former energy range is concerned (see Fig. 3a), the K-edge EELS spectra acquired from the nanostructure before thermal annealing are entirely lacking the articulated post-edge structures observed in the corresponding spectra acquired from the surrounding diamond matrix. Remarkably, this clear distinction in EELS spectral features is fully preserved after the annealing step, thus indicating that no phases attributable to crystalline diamond can be detected upon thermal processing.
Fig. 3 Results of cross-sectional EELS spectroscopy of the nanostructures. K-edge absorption features of both the nanostructure (red plots) and the surrounding diamond matrix (blue plots) are reported for both the as-implanted (a) and thermally processed (b) sample. The pre-edge peak at ∼285 eV, which is attributed to π-bonded carbon, is visible in a, while it is entirely absent in (b). (c) Present experimental data are compared to the data reported for fully-sp3 bonded amorphous carbon in ref. 20, as well as to the characteristic spectra of glassy carbon (in which the ∼285 eV feature is labeled as *) and nanocrystalline diamond (whose characteristic post-edge structure is labeled as §). Low-loss spectra exhibit a characteristic downshift in the plasmon-related features, both before (d) and after (e) thermal processing. |
Note that the spectra acquired from the nanostructures before thermal treatment exhibit a well-defined (although not particularly intense) absorption pre-edge peak at ∼285 eV that is unequivocally attributed to π-bonded carbon,33 thus indicating that a fraction of sp2 bonds is indeed present in the as-implanted phase. Contrarily to what is commonly observed in amorphized carbon, this spectral feature does not increase upon thermal annealing, but rather completely disappears, which unequivocally indicates that the sp2 bonds are absent from the annealed nanostructure within the detection limit of this very sensitive technique.
For the sake of comparison, Fig. 3c reports our experimental data together with the EELS spectrum collected from the fully sp3-bonded amorphous carbon phase investigated in ref. 20: the mutual similarity is striking, particularly considering that both spectra entirely lack the features associated with glassy carbon and nanocrystalline diamond (marked as * and § in the respective reference spectra). The EELS features in the low-energy-loss range reported in Fig. 3d and e exhibit plasmon peaks which are indicative of the electron densities in the corresponding phases. Both before and after thermal annealing, a systematic shift to lower energy losses of the plasmon peak positions is observed from the nanosctructures with respect to the ∼34 eV peak, which is characteristic of the surrounding diamond matrix.14,34–36 These shifts can be interpreted on the basis of the lower atomic density of the nanostructures, by adopting the following formula:36–38
Given the mass density (i.e. 3.515 g cm−3) and plasmon peak position (i.e. ∼34 eV) of the surrounding diamond matrix, and under the assumption that the same electron effective mass can be adopted for the different phases under investigation,38 it is possible to estimate the electron and (and thus mass) density within the nanochannels from the position of the corresponding plasmon peaks. Under these approximations, the (31.0 ± 0.3) eV and (32.6 ± 0.3) eV plasmon peak positions measured from the nanostructures before and after thermal annealing yield mass density estimations of (2.92 ± 0.06) g cm−3 and (3.27 ± 0.07) g cm−3, respectively.
This result is indicative of the fact that: (i) the implantation process results in a substantial density variation within the nanochannels, despite the strong compressive stress exerted by the rigid surrounding diamond matrix; and (ii) the disappearing of sp2 bonds within the nanochannels upon thermal annealing determines a substantial increase of mass density with respect to the as-implanted sample, which closely approaches the density of pristine diamond and is fully compatible with the estimation (i.e. 3.3 g cm−3) provided for the fully sp3-bonded amorphous carbon phase reported in ref. 20.
Finally, we remark that thermal annealing results in radically different structural features across the previously defined “continuous shallow layer” located above the nanochannels (see Fig. 4a). While (as much as observed from the nanochannels) the continuous shallow layer is characterized by TEM diffractometry features that are indicative of a fully amorphized phase (see Fig. 4b), the EELS spectrum (see Fig. 4c) exhibits the strong absorption pre-edge peak at ∼285 eV. Such a difference is attributed to the different geometries of the two types of structures, while all other fabrication and processing parameters (irradiation, thermal annealing) are the same. This strongly indicates that the peculiar stress field developed in correspondence of the nanostructures is primarily responsible for the formation of an amorphous full-sp3 network.
Fig. 4 Structural properties of the continuous shallow layer located above the nanostructures, following thermal annealing. (a) Cross sectional bright-field TEM micrograph from the annealed structure: the labels indicate the diamond cap layer, the continuous shallow layer and the nanochannel embedded in the bulk diamond; the inset labeled by “SRIM” reports the two-dimensional damage density plot reported in Fig. 1c, for sake of comparison. (b) TEM diffraction pattern collected from a random region of the continuous shallow layer. (c) Corresponding EELS spectrum, clearly exhibiting the strong absorption pre-edge peak at ∼285 eV which is indicative of a large fraction of sp2 bonds. |
Our interpretation is confirmed by 2D finite element method (FEM) mechanical analysis. By simulating the constrained expansion undergone by the two implanted regions, it is possible to highlight a significant volumetric stress build up in the 40–50 GPa range (consistently with results in ref. 20) upon ion implantation.
However, as shown in Fig. 5a, the shallow layer (region encircled by the dashed line) does not undergo significant stresses in the vertical (y) direction, while the latter tends to accumulate in the end-of-range region of the nanochannel, due to the confining effect of the surrounding pristine diamond material. The depth variation of σy stresses in the nanochannel (Fig. 5b), which reach a peak value of about 48 GPa, while remaining negligible in the shallow layer (shaded area), are identified as the factors that are responsible for the different structures observed in the two regions. Fig. 5c and d show the results of analogous numerical simulations carried for the nanostructure after thermal annealing, under the assumption that the latter process results in a conversion to a graphitic phase. In this case, a distribution of relatively high compressive stresses (in the GPa range) still persists. It is worth remarking that stress fields of this order of magnitude are indeed observed in graphitized microstructures created with this technique, such as the ones reported in ref. 22. The experimental observation that thermal annealing does not result in the graphitization of the nanostructure is therefore attributed to the fact that the strong stress fields established around the nanostructure upon ion implantation are instead fully relaxed (i.e. residual stresses in the Pa range) upon the formation of a fully sp3-bonded amorphous carbon phase, as confirmed by the simulations reported in Fig. 5e and f. In our interpretation, the stress state responsible for the transition to sp3 bonds is established before the thermal annealing. The initially strongly stressed nanochannels therefore transition to a fully sp3-bonded amorphous phase upon thermal annealing, and only subsequently are the stresses relaxed.
As mentioned above, an experimental assessment of the local stresses established across amorphized/graphitized structures embedded in the diamond matrix is in general possible by means of confocal micro-Raman spectroscopy, both before and after thermal annealing, as already demonstrated for micrometer-sized regions.22 This was not possible (either by conventional or tip-enhanced Raman spectroscopy) in the case of the nanostructures reported in the present work, due to (respectively) limited spatial resolution and spectral sensitivity of the available techniques. Nonetheless, it is worth remarking that EELS provided direct experimental insight into the local mass density of the nanostructures, which directly translated into the above-described numerical simulations.
The results presented in this work provide important information regarding the mechanisms leading to the formation of fully-sp3-based amorphous phases in carbon, and display appealing applications in fields where high-pressure carbonaceous phases could be implemented in integrated devices, such as for example room-temperature superconducting devices.39
Besides the fundamental relevance of this finding in the understanding of the formation of exotic carbon phases, the use of focused/collimated ion beams enables the direct fabrication of such nanostructures in complex patterns and arrangements, with appealing perspectives in nanomechanical systems and integrated nano-optics.
To gain further insight in the process of amorphization and estimate the stresses acting on the irradiated region, 2-D FEM simulations were performed using Comsol Multiphysics. Consistently with the results of cross-sectional TEM microcopy (see Fig. 4a), a 5 × 5 μm2 diamond cross section was considered, in which a 0.1 × 1.8 μm2 implanted diamond strip is incorporated, representing the nanochannel, including a terminal trapezoidal region to account for straggling effects, and a 5 × 0.5 μm2 strip representing the shallow layer (see Fig. 5). For the as-implanted sample, material parameters are: diamond density ρd = 3.52 g cm−3, amorphous carbon density ρaC = 2.06 g cm−3, diamond Young's modulus Ed = 1220 GPa, amorphous carbon Young's modulus EaC = 21.38 GPa.42,43 The density of the implanted region before annealing was calculated as a function of the induced damage density as reported in ref. 44, i.e.:
ρ(y) = ρd − (ρd − ρaC)Φ(y) |
The Young's modulus decrease as a function of the vacancy density can be expressed as:
E(y) = Ed − (Ed − EaC)Φ(y) |
The density decrease due to irradiation generates a constrained expansion of the implanted volume, i.e. residual strains in the i = x,y directions that can be expressed as follows:
In the case of the sample after thermal annealing, the simulations were carried under two alternative hypotheses, i.e. a conversion of the nanostructure to a graphitic phase (see Fig. 5c and d) or to a fully sp3-bonded amorphous carbon phase (see Fig. 5e and f). In the former case, the following structural parameters were adopted: ρg = 2.1 g cm3, Young's modulus Eg = 21 GPa. In the latter case, the following structural parameters were adopted: ρa–d = 3.3 g cm3 [ref. 20, this work], Ea–d = 1123 GPa.20
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