We report on fundamental discovery of conversion of amorphous carbon into diamond by irradiating amorphous carbon films with nanosecond lasers at room-temperature in air at atmospheric pressure. We can create diamond in the form of nanodiamond (size range <100 nm) and microdiamond (>100 nm). Nanosecond laser pulses are used to melt amorphous diamondlike carbon and create a highly undercooled state, from which various forms of diamond can be formed upon cooling. The quenching from the super undercooled state results in nucleation of nanodiamond. It is found that microdiamonds grow out of highly undercooled state of carbon, with nanodiamond acting as seed crystals.

The carbon to diamond conversion at ambient pressures and lower temperatures is scientifically challenging with immense technological significance.1–3 Conversion of carbon, one of the most abundant materials in the Earth’s crust, into most precious material diamond has been a cherished goal of the scientists all over the world for the longest time. Diamond is one of the most desirable materials with many applications ranging from abrasives, protective coatings, and biomedical applications to superior diamond electronics, photonics, and display devices. Conventional bulk processing involves high pressures and temperatures,1 and chemical vapor deposition for thin films requires high temperatures in the presence of hydrogen.4 These requirements lead to low production volumes and high costs. Conversion of SiC into nanodiamond has been reported at temperatures ∼1000 °C under flowing hydrogen and chlorine gases at ambient pressures.5 Here, we show that a direct conversion of carbon into diamond can occur in air at ambient temperatures and pressures without any need for catalysts and hydrogen to stabilize sp3 diamond bonding.

According to the equilibrium (P vs. T) phase diagram (shown in Fig. 1 with an inset at low pressures), graphite, diamond, liquid, and vapor are thermodynamically stable forms of carbon. At low pressures, graphite converts into vapor above around 4000 K. According to this phase diagram, diamond synthesis from liquid carbon will require even higher temperatures and pressures as the graphite/diamond/liquid carbon triple point occurs at 5000 K/12 GPa, where 1 GPa = 9869 atm. Consistent with the phase diagram, diamond can exist in the interiors of the outer planets (Uranus and Neptune) and Earth’s mantle, where pressure/temperature is 600 GPa/7000 K and 135 GPa/3500 K, respectively. Currently, diamond powders are synthesized by graphite to diamond conversion at high pressures and temperatures. Graphite can be transformed into diamond above about 2000 K at 6-10 GPa using liquid metal catalysts which are used for commercial synthesis of diamond.1 Because of the high binding and activation energy of transformation, carbon polymorphs can exist metastably well into a P-T region, where a different phase is thermodynamically stable. As an example, diamond survives indefinitely at room temperature, where graphite is the stable form. The vertical straight line in the phase diagram near room-temperature is based upon the results from room-temperature compression of graphite. Room-temperature compression of graphite along the c-axis shows the development of sp3 bonding at about 12-14 GPa, which decreases electrical conductivity and optical reflectivity and increases optical transmittance by removing conduction electrons. The decompression from 12 to 14 GPa state reverts the phase transformation to graphite phase; however, the transformation is irreversible at higher pressures and upon heating, this phase transforms into very hard hexagonal diamond.1 

In the equilibrium phase diagram of carbon, we introduce an amorphous metastable phase of carbon, where bonding characteristics are a mixture of graphite (sp2 bonded) and diamond (sp3 bonded). By introducing metastable phase of diamondlike carbon, is it possible to convert this phase into diamond at lower temperatures and ambient pressures? We show that this amorphous phase can be melted at a much lower temperature (over 1000 K less than the melting point of crystalline carbon) to create a highly undercooled state of carbon liquid. We believe that diamond nanocrystallites are nucleated from four-fold coordinated (sp3) bonded carbon present in this super undercooled state of quenched carbon (Q-carbon). We find that amorphous phase of carbon is critical for the large undercooling needed for the formation of diamond phase. In the case of amorphous Ge and Si, highly undercooled states have been found to occur at 241 K and 336 K below their respective melting points.6,7 By scaling with the melting point of carbon, we estimate the undercooling in carbon to be as high as 1000 K. This undercooling can shift amorphous carbon/diamond/liquid carbon triple point to 4000 K or lower at ambient pressures. The interface temperature was estimated to be 4000 K by laser-solid interactions calculations, Simulation of Laser Interaction with Materials (SLIM) model developed by Singh and Narayan.6 This is a rather drastic change from graphite/diamond/liquid carbon triple point at 5000 K/12 GPa. At these transition temperatures, Gibbs free energy of amorphous carbon equals the free energy of highly undercooled liquid and metastable diamond phase which is quenched and retained at room temperature. We show that the creation of this super undercooled state is critical to the nucleation of nanodiamond crystals which provide as a seed for the growth of microdiamond crystals. In addition to undercooling, the bonding characteristics of amorphous carbon play an important role in the nucleation of diamond phase. It should be pointed out that graphite is sp2 bonded and diamond is sp3 bonded, whereas amorphous carbon films have a mixture of the two bonding states. An optimum ratio of sp3 to sp2 bonding seems to be necessary for the nucleation of diamond nanocrystallites.

FIG. 1.

Carbon phase diagram (P vs. T) following Bundy et al.1 which has amorphous diamondlike carbon melting at 4000 K at ambient pressures (dotted green line).

FIG. 1.

Carbon phase diagram (P vs. T) following Bundy et al.1 which has amorphous diamondlike carbon melting at 4000 K at ambient pressures (dotted green line).

Close modal

The phenomenon of undercooling in amorphous Ge and Si was one of the hottest topics at the early Materials Research Society (MRS) meetings7–10 in the context of pulsed laser annealing,11 solute trapping,12 and formation of metastable supersaturated semiconductor alloys.13 However, undercooling in amorphous carbon was not pursued with the same vigor, as carbon/graphite is supposed turn into vapor at ambient temperatures and pressures.1 Previous works on nanosecond laser melting of carbon implanted copper showed evidence of diamond formation. In this case, carbon was implanted into Cu(100) single crystals to a high dose, followed by laser melting 3-4 J cm−2 using XeCl excimer laser (wavelength 308 nm and pulse duration 45 ns). This melting zone refined all the Cu in the melted region to the surface and diamond nucleated on Cu substrate which is lattice matched with diamond.2,3 However, the undercooling needed for diamond nucleation was harder to achieve over a large area consistently. Related studies on laser melting of 73Ge+ and 75As+ implanted highly oriented pyrolytic graphite (HOPG) using a ruby laser (wavelength 693 nm, pulse duration 30 ns, energy density 0.6-3.0 J cm−2) did not reveal any evidence of formation of diamond phase, presumably because of highly conducting epitaxial graphite substrate.14,15 The rapid heat diffusion did not produce sufficiently high undercooling and time needed for diamond nucleation. In addition, no sp3 bonded carbon was present in the molten carbon to provide a seed for diamond nucleation. More recent experiments on melting on HOPG using 1-ms pulses Nd:YAG (wavelength 1060 nm and power 10 kW) under the He gas pressure of 12 MPa have reported occasional presence of a very small amount of cubic phase of diamond using TEM electron diffraction. No Raman spectra were reported from these experiments perhaps due to very small amount of diamond phase.16,17 It should be pointed out that undercooling with 1-ms pulses of Nd:YAG on HOPG substrates is expected to be even smaller than nanosecond laser pulses.

The primary focus of this paper is on direct conversion of carbon into diamond by nanosecond laser melting of diamondlike amorphous carbon films on sapphire and glass substrates, where much higher undercooling up to 1000 K can be achieved by ArF excimer laser (wavelength 193 nm and pulse duration 20 ns). These undercooling values for amorphous carbon are considerably higher than those achieved during melting of crystalline carbon such as HOPG samples. It should be emphasized that the presence of certain sp3 fraction and undercooled melt lifetime is critical to diamond nucleation. The nanodiamonds nucleate from the undercooled state, which provide a seed of microdiamond formation. We also discuss the role of sp3 bonded carbon in diamondlike films on the nucleation of diamond nanocrystallites.

The amorphous carbon films containing both sp2 and sp3 bonding states were deposited on sapphire (c-plane) and glass substrates by using KrF laser (pulse duration = 25 ns, wavelength = 248 nm, energy density = 3.0 J cm−2) to a thickness of 50-500 nm. These films were characterized by TEM and Raman and found to be amorphous containing Raman signature on sapphire (Diamondlike Carbon (DLC) broad peak = 1580 cm−1) with estimated sp3 fraction varying from 20% to 50%. The Raman spectra for films on glass substrates, contained D (1349 cm−1) and G (1580 cm−1) peaks with considerably less sp3 around 20%-25%. The films on sapphire contained a single broad peak centered on 1580 cm−1 with sp3 fraction over 40%. These films were irradiated in air with ArF laser pulses (pulse duration = 20 ns, wavelength = 193 nm, energy density = 0.3-0.6 J cm−2). High-resolution SEM and electron backscattered diffraction (EBSD) with characteristic Kikuchi patterns were carried out using FEI Verios 460L SEM and FEI Quanta 3D FEG FIB-SEM for phase identification and determination of grain orientation. Aberration corrected STEM-FEI Titan 80-300 and JEOL-2010 STEM/TEM were used electron energy loss spectroscopy (EELS) with resolution of 0.15 eV and high resolution TEM (point-to- point TEM resolution 0.18 nm; STEM-Z resolution 0.08 nm with information limit of 0.06 nm).

Fig. 2(a) shows the formation of nanodiamond (10-50 nm) on a sapphire substrate after a single laser pulse of 0.5 J cm−2. At a higher energy density (0.6 J cm−2), microdiamond (300-800 nm) is formed, covering the entire area of the substrate (Fig. 2(b)). It is interesting to note that some of these microdiamond crystallites are surrounded by high-index planes, as determined by EBSD analysis, and are not as faceted with low energy planes, as found typically in CVD diamond films. The crystallites with high-index planes have been shown to be catalytically more active than low-index ones. The mechanism of formation of microdiamond is shown in Fig. 2(c), where it is growing out from a nanodiamond formed from the undercooled carbon in the form of a filament.

FIG. 2.

High resolution SEM micrographs: (a) mixture of nanodiamond and microdiamond; (b) entire area covered with microdiamond; and (c) nucleation of microdiamond from nanodiamond filaments which were quenched after the formation at 4000 K.

FIG. 2.

High resolution SEM micrographs: (a) mixture of nanodiamond and microdiamond; (b) entire area covered with microdiamond; and (c) nucleation of microdiamond from nanodiamond filaments which were quenched after the formation at 4000 K.

Close modal

Fig. 3(a) shows Raman results from sapphire substrate after a single laser pulse of ArF laser (energy density 0.6 J cm−2). A sharp diamond peak at 1331.3 cm−1 along with sapphire peaks S1 and S2 (at 1360 and 1375 cm−1) and small G peak of residual unconverted amorphous graphite are observed. In this case, as-deposited films had a sp3 fraction, closer to 35%. The Raman spectrum from the filament (shown in Fig. 2(c)) contained a diamond peak at 1331 cm−1 with larger FWHM along with sapphire and the G peak. A slight down shift of the primary Raman peak and a bump at 1140 cm−1 in the filament spectrum are characteristic of sp2 bonded carbon at the grain boundaries in nanodiamond. Fig. 3(b) shows Raman spectra before and after laser annealing of diamondlike carbon films on a glass substrate. As deposited amorphous carbon films show both D (1350 cm−1) and G (1580 cm−1) peaks with estimated sp3 fraction about 25%. After a single laser pulse of ArF laser energy density (0.6 J cm−2), sharp diamond peaks at two different places on the same sample at 1333.6 cm−1 are observed. A slight upshift in the diamond Raman may be related to the residual compressive stresses and defects in diamond thin films. The upshift (Δω) is related to Δω (in cm−1) =2.2 ± 0.10 cm−1 GPa−1 along the [111] direction, Δω (in cm−1) =0.73 ± 0.20 cm−1 GPa−1 along the [100] direction, and Δω (in cm−1) =3.2 ± 0.23 cm−1 GPa−1 for the hydrostatic component.18 

FIG. 3.

Raman spectra from films on sapphire and glass substrates: (a) from sapphire substrate on diamond Raman peak at 1331.3 cm−1, on nanodiamond filaments and as-deposited films. S1 and S2 peaks belong to sapphire and D and G peaks to diamondlike carbon; (b) from glass substrate on diamond Raman peak at 1333.6 cm−1 from two different areas and D and G peaks for diamondlike carbon.

FIG. 3.

Raman spectra from films on sapphire and glass substrates: (a) from sapphire substrate on diamond Raman peak at 1331.3 cm−1, on nanodiamond filaments and as-deposited films. S1 and S2 peaks belong to sapphire and D and G peaks to diamondlike carbon; (b) from glass substrate on diamond Raman peak at 1333.6 cm−1 from two different areas and D and G peaks for diamondlike carbon.

Close modal

Fig. 4(a) is a cross section image of the sample using high-resolution SEM, where the growth of microdiamond on sapphire substrate is shown in detail. HRTEM image from a diamond microcrystallite from this area is shown in Fig. 4(b), where the 110 cross section has two sets of {111} planes of diamond clearly imaged. The characteristic 110 diamond selected-area-diffraction pattern is included in the inset. These films have shown 110 diamond texture, which is consistent with 110 texture of silicon formed under rapid unseeded crystallization (18). Fig. 4(c) shows cross section TEM image from where EELS was collected to study sp3*) bonding characteristics of diamond. The Fig. 4(c) also reveals the presence of nanodiamond layer (marked by an arrow) adjacent to sapphire substrate, which provides a seed for microdiamond formation, as shown in Fig. 2(c). Using Titan HRTEM/STEM, EELS with a resolution of 0.15 eV is shown in Fig. 4(d). The spectrum contains the sharp edge at 288 eV with a peak at 292 eV, corresponding to sp3*) bonding, which is a signature EELS spectrum for diamond.

FIG. 4.

High-resolution TEM: (a) SEM image of cross section TEM sample with diamond on sapphire; (b) 110 HRTEM image showing with resolution of 0.18 nm with inset 110 electron diffraction pattern; (c) cross section from the same sample showing filament (shown by red arrows) near the sapphire substrate and microdiamond; and (d) diamond EELS spectrum from microdiamond in (c) using FEI Titan with resolution of 0.15 eV.

FIG. 4.

High-resolution TEM: (a) SEM image of cross section TEM sample with diamond on sapphire; (b) 110 HRTEM image showing with resolution of 0.18 nm with inset 110 electron diffraction pattern; (c) cross section from the same sample showing filament (shown by red arrows) near the sapphire substrate and microdiamond; and (d) diamond EELS spectrum from microdiamond in (c) using FEI Titan with resolution of 0.15 eV.

Close modal

The diamond structure determination and phase identification have been also carried out using EBSD by using field-emission scanning electron microscope. The results are shown in Fig. 5(a) from a microdiamond with SEM micrograph, characteristic diamond Kikuchi pattern, and orientation relationship of the diamond with respect to the substrate normal. The SEM micrograph shows needle shaped microdiamond growing from the nanodiamond layer near the sapphire substrate, as shown in Fig. 4(c). Fig. 5(b) shows similar EBSD results from a flat area of the diamond film.

FIG. 5.

Electron Backscatter Diffraction (EBSD) patterns; (a) 110 EBSD Kikuchi pattern from the encircled nanodiamond needle; (b) EBSD from a flat area encircled.

FIG. 5.

Electron Backscatter Diffraction (EBSD) patterns; (a) 110 EBSD Kikuchi pattern from the encircled nanodiamond needle; (b) EBSD from a flat area encircled.

Close modal

The phenomenon of undercooling and formation of nanocrystalline and microcrystalline silicon has been studied in detail in the model system of amorphous silicon, and there are interesting parallels here for the formation of nano- and microdiamonds from amorphous carbon.18–20 Fig. 6(a) illustrates this super undercooling phenomenon in amorphous silicon, which had 500 nm thick amorphous silicon layer after (100) silicon was implanted with 200 keV self-ions to a dose 1.5 × 1016 cm−2. Here, after a single pulse of nanosecond laser (wavelength 693 nm, pulse duration 12 ns, energy density 0.8 J cm−2), we observe 100 nm remaining (amorphous silicon), which is followed by 100 nm nanocrystalline silicon and 300 nm microcrystalline silicon. In the diffraction from the nanocrystalline region (Fig. 6(b)), there was no evidence for the formation of amorphous phase. The formation of 100 nm wide nanocrystalline region occurs by nucleation in the super undercooled state of silicon.18–20 The detailed modeling showed the nucleation of nanocrystallite (10 nm size) silicon from super undercooled melt and the time-resolved reflectivity measurements showed the importance of laser and substrate parameters, particularly the thermal conductivity of residual underlying amorphous layers, in achieving the undercooled state of silicon.20,21 These results of undercooling in silicon emphasize the importance of low thermal conductivity substrates of sapphire and glass in the case of our diamond experiment. The formation of nanocrystalline region was found to be not a function of dopants, indicating the importance of homogeneous nucleation. In the case of amorphous Ge and Si, highly under cooled states have been found to occur at 241 K and 336 K below their respective melting points.6–10 Under the super undercooled state of silicon, bulk nucleation of nanocrystalline silicon was directly observed, where the nanocrystallites provided seed for macrograined polycrystalline silicon, as shown in TEM micrograph of Fig. 6. From the distribution of impurities, the solidification in the microregion was estimated to be 3-5 ms−1, while in the nanoregion, the solidification velocity approached 10-12 ms−1.18–21 

FIG. 6.

Super undercooling in the model system of amorphous silicon: (a) cross section TEM image showing formation of nano- and microcrystalline regions; (b) corresponding selected area diffraction pattern showing characteristic rings.

FIG. 6.

Super undercooling in the model system of amorphous silicon: (a) cross section TEM image showing formation of nano- and microcrystalline regions; (b) corresponding selected area diffraction pattern showing characteristic rings.

Close modal
The formation of nanodiamond occurs as a result of nucleation of diamond phase from the highly undercooled state of pure carbon. The diamond formation occurs by a homogenous nucleation from highly undercooled state of pure carbon, as shown in Fig. 7, where diamond nuclei vary from 2 nm to 8 nm. The purity was assessed by EELS, Energy-dispersive X-ray spectroscopy (EDAX), and SIMS analyses and impurity content was found to be less than 1-10 ppb. The Gibbs free energy of diamond nuclei (ΔGT) consists of gain in volume energy (ΔGV) and expense of surface free energy (ΔGS),
Δ G T = Δ G V + Δ G S ,
(1)
Δ G T = 4 3 π r 3 ρ M m Δ H m T m Δ T u + 4 π r 2 r s ,
(2)
where “r” is the radius of diamond nucleus, ( ρ M m Δ H m T m Δ T u ) is the gain in free energy for the formation of diamond nucleus from the undercooled state, ΔTu is the undercooling from Tm to Tr (temperature of nucleation), ρ is the solid diamond density, Mm is the molar mass, ΔHm is the latent heat of melting, Tm is the melting point of carbon, and rs is the surface free energy between diamond nuclei and carbon liquid.
FIG. 7.

Homogeneous nucleation of diamond (2-8 nm diameter) after laser irradiation at energy density of 0.55 J cm−2.

FIG. 7.

Homogeneous nucleation of diamond (2-8 nm diameter) after laser irradiation at energy density of 0.55 J cm−2.

Close modal
The maximum of ΔGT corresponds to the diamond reaction barrier at a critical size of r, where
r * = 2 r s T m M m Δ H m Δ T u ρ ,
(3)
Δ G T * = 16 π r s 3 T m 2 M m 2 3 Δ H m 2 Δ T u 2 ρ 2 .
(4)
Rate of nucleation (I) is governed by
I = A exp Δ G T * K T r ,
(5)
where Tr = Tm − ΔTu, A = n(kT/h) exp (−ΔFA/kT), I = number of diamond nuclei cm−3 s−1, n = number density of atoms, and ΔFA is free energy of activation across the liquid-solid interface. Our calculated values for 5 and 10 Å for critical size of diamond crystallite lead to A values of 1025 and 1024 cm−3 s−1, respectively, which are consistent with our earlier results on homogeneous nucleation in super undercooled amorphous silicon20 and theoretical estimates in carbon.22 At the formation temperature, Tr, the free energy of metastable diamond, highly undercooled carbon liquid, and amorphous diamond like carbon become equal.

Diamond nucleation is difficult due to large rs but the large undercooling ΔTu drives the r* lower and helps the nucleation. Also, ΔG is lowered considerably under large undercooling due to ΔTu−2 dependence. The lower ΔG also enhances nucleation rate of nanodiamond from the highly undercooled state and these nanodiamonds provide as a seed for microdiamond growth. It should be emphasized that the undercooling should be retained for a sufficient time for diamond nanocrystallite not only able to nucleate but also grow. For a 10 nm size nanodiamond, the estimated growth time is about 5-10 ns. This time requirement emphasizes the importance of thermal conductivity of the substrate during rapid pulse laser heating. A rough estimate of r* for diamond nuclei from Equation (3) is 5-10 Å, where diamond surface free energy rs is 0.6 J m−2, Tm = 4000 K, ΔHm = 1.0 eV/atom, ΔTu = 500-1000 K, and ρ = 3.5 gm/cm−3. Thus, diamond nucleation from highly super undercooled carbon can occur readily.

This transformation from super undercooled state of carbon into diamond, which can occur at 4000 K and ambient pressures, modifies the equilibrium carbon phase diagram, as shown in Fig. 1. The extension (shown as dotted green line) of the phase diagram (P vs. T) is based upon Simon equation, where P = P0 + a[(T/Tr)c − 1], where we used P0 = 10−4 GPa at Tr = 4000 K from our experiments and the experimental data of Bundy et al.1 to estimate a and c parameters. From the Simon equation, (dP/dT)Tr = ac/Tr was determined and the results compared with the value estimated from Clausius-Clapeyron equation, where (dP/dT)Tr = ΔHm/(TrΔV), where ΔV is the change in volume from super undercooled carbon into diamond.

Undercooling during nanosecond laser (KrF laser λ = 0.248 μm) have been measured in crystalline Ge and Si using nanosecond resolution time-resolved x-ray diffraction in a Synchrotron. These results showed undercooling as high as 110 ± 30 K on (111) Si during ∼10 ms−1 melting and ∼6 ms−1 regrowth velocity. Thus, these undercooling values in crystalline Ge and Si are roughly half that predicted theoretically for amorphous Ge and Si. The undercooling was also found to be orientation and regrowth velocity dependent: ∼5 K/m s on (100)Ge and ∼20 K/m s on (111)Ge.23,24

Basharin et al.16 have estimated the temperature dependence of Gibbs free energy of graphite, metastable liquid carbon and diamond at a low pressure of 120 MPa of helium and found that free energy of liquid carbon can be equal to that of diamond at 4160 ± 50 K, which is considerably lower than melting point of liquid graphite, at a pressure of 120 MPa of helium. Thus, it is possible that diamond can be nucleated from super undercooled state at 4160 K. However, Basharin et al.17 tried to quench diamond using 1-ms pulsed laser melting (wavelength of λ = 1.06 μm and a power of 10 kW) of HOPG graphite with limited success due to lower undercooling with 1 ms lasers on a crystalline graphite substrate.16,17

Bonding and structure of liquid carbon have been studied using time-resolved x-ray absorption spectroscopy25 and results compared with first-principles molecular dynamics simulations.26 These results suggest that local bonding structure of the liquid can be varied from a mixture of twofold and threefold coordinated atoms at low density (1.27 g cm−3) to fourfold coordinated atoms at high density (3.02 g cm−3). The latter liquid structure corresponds to liquid carbon structure at high pressures. Thus, the structure of amorphous carbon (containing >30%–40% sp3 bonding) at low pressures under super undercooling is similar to the structure of liquid carbon at high pressures, where conversion into diamond phase occurs.

In summary, our results clearly show that diamond can be formed at ambient temperatures and atmospheric pressure in air from super undercooled state of carbon without any need for any catalysts and hydrogen to stabilize sp3 diamond bonding. Thus, amorphous state of carbon, laser parameters, and film substrate characteristics determines the temperature distribution and undercooling and plays a critical role in nucleation and growth of diamond. By using nanosecond pulsed lasers, we are able to achieve super undercooled state of carbon, which can be quenched to form nanodiamonds from which microdiamonds can grow out with a large yield. Extent of undercooling, sp3 content of amorphous carbon, laser and substrate parameters, and melt life-time available for nucleation and growth of diamond from the liquid phase, all play a critical role in the conversion of carbon into diamond at ambient temperatures and pressures (atmospheric) in air. Since nanosecond laser heating and temperature distributions are confined spatially and temporally in such a way that while surface layers melt, the underlying substrates stay close to the ambient. Thus, this method can be used to create epitaxial diamond nanostructures by providing appropriate template and deposit diamond on heat sensitive substrates such as polymers and plastics. This method provides a rather inexpensive and useful way to convert carbon into diamond at reduced temperatures and ambient pressures in air. This discovery opens a new chapter in synthesis and processing of nanodiamond and microdiamond for a variety of applications ranging from abrasive powders, novel catalytic properties, smart displays, myriads of biomedical and microelectronic, and nanoelectronic applications.

We are grateful to Fan Family Foundation Distinguished Chair Endowment for J. Narayan. We are also very pleased to acknowledge technical help and useful discussions with John Prater, Jim LeBeau, Weizong Xu, Xiahan Sang, Roger Narayan, and Sandhyarani Punugupati.

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