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oastm_01
The number of grains evaluated in each micrograph was about 380 except for the totally recrystallized sample annealed at 673K, for which about 130 grains were considered.
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oastm_01
train5
steel
oastm_01
The error in the 1% range is due to uncertainties of assigning contrast features to grain boundaries.
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oastm_01
train5
steel
oastm_01
Results
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oastm_01
train5
steel
oastm_01
Correlation between microstructure and free volumes
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6,832
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oastm_01
train5
steel
oastm_01
Fig. 2 shows a typical dilatometric length contraction Δl/l0, indicating the annealing out of defects associated with the release of free volume upon linear heating of the HPT-deformed Cu sample.
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oastm_01
train5
steel
oastm_01
Three substages, A, B and C, can clearly be discerned.
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oastm_01
train5
steel
oastm_01
In order to correlate these substages with annealing of specific types of lattice defects, a microstructural characterization by means of scanning electron microscopy was performed at the onset and at the end of each substage, as indicated by arrows in Fig. 2.
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oastm_01
train5
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oastm_01
In stage A, only a minor increase in the crystallite size from 209±4nm (as-deformed, Fig. 3a) to 222±4nm (end of stage A, Ta=413K,5K/min-1, Fig. 3b) is observed.
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oastm_01
train5
steel
oastm_01
Substantial crystallite growth up to 764±4nm (Ta=468K, Fig. 3c) takes place in the subsequent distinct annealing stage B.
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oastm_01
train5
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oastm_01
The broad annealing stage C is accompanied by further crystallite growth up to 12.09±0.04μm (Ta=673K, Fig. 3d).
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oastm_01
train5
steel
oastm_01
In contrast to HPT-deformed Ni (see below), both the crystallite shape and the dilatometric length change are isotropic.
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oastm_01
train5
steel
oastm_01
Quite similar dilatometric curves were obtained for the three measuring directions (axial, tangential and radial, Fig. 1).
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oastm_01
train5
steel
oastm_01
From the measurements of nine dilatometric samples, a total length change Δl/ltotal=(7.02±1.92)×10-4 is deduced that comprises a fraction (Δl/l0)Stage A=(1.32±0.46)×10-4 for stage A, (Δl/l0)Stage B=(1.85±0.37)×10-4 for stage B and (Δl/l0)Stage C=(3.77±1.33)×10-4 for stage C.
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oastm_01
train5
steel
oastm_01
The results and uncertainties represent the mean value and standard deviation of the nine dilatometric measurements.
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Since, during annealing in stage A, the crystallite size and, therefore, the fraction of grain boundaries remain nearly constant, the length contraction in this stage has to be attributed to the annealing out of crystal lattice defects as well as to a relaxation of grain boundaries (see Section 4).
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In the clearly defined subsequent stage B, the length change is predominantly caused by the removal of relaxed grain boundaries, as evidenced by the increase in the crystallite size by more than a factor of 3 in this narrow temperature regime.
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This allows one to determine the grain boundary excess volume eGB from the relative length change (Δl/l0)Stage B in this stage according to(1)Δll0Stage B=eGB1dini-1dfinwhere dini=222nm and dfin=764nm denote the crystallite diameter at the beginning and at the end of stage B, respectively.
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From the mean value (Δl/l0)Stage B=(1.85±0.37)×10-4 of nine samples, an apparent grain boundary excess volume eGB′ of (0.58±0.13)×10-10m can be derived.
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However, for a more detailed analysis, it should be taken into account that the onset of the annealing processes of stage C already occurs during stage B.
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Subtracting this contribution by means of extrapolation of stage C to lower temperatures into the stage B regime yields a reduced effective value (Δl/l0)Stage B, eff=(1.48±0.29)×10-4 associated with the removal of grain boundaries, resulting in a corrected value of the grain boundary excess volume of eGB=(0.46±0.11)×10-10m.
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Kinetics and comparison with DSC
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The shift of the dilatometric length change in stage B with heating rate (see the inset in Fig. 2) provides insight into the kinetics of crystallite growth.
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Analyzing this shift for the nine dilatometric measuring runs according to Kissinger (for details of the method see Ref.
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[15]) yields an activation energy Q of crystallite growth of 0.99±0.11eV (see the Kissinger plot, Fig. 4).
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For the sake of comparison, a series of 14 DSC measurements with different heating rates were taken on samples prepared from the same HPT disk.
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As shown in Fig. 5, the dilatometric stage B due to crystallite growth is accompanied by a pronounced heat release.
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Kissinger analysis of the shift of this DSC peak with heating rate yields an activation energy Q of 0.96±0.06eV, which is in excellent agreement with the dilatometric measurements (see Fig. 5).
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From the heat release upon linear heating through stage B, a mean value of the enthalpy of ΔH=-0.92±0.06Jg-1 for the exothermic process is deduced from the various measuring runs.
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Attributing this enthalpy release exclusively to the removal of grain boundaries in stage B, a specific grain boundary energy(2)γ=Hρ3dini-1-dfin-1=0.85±0.08Jm-2is estimated using the initial and final crystallite diameters of stage B, as given above (dini=222nm, dfin=764nm), as well as the Cu bulk value of 8.92gcm-3 for the mass density ρ.
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This value of γ is typical of relaxed grain boundaries in Cu (see Ref.
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[16] and references therein).
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The results for Cu of the present work are summarized in Table 1, together with data previously obtained for Ni [9,15] for comparison.
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Discussion
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For the discussion of the present dilatometric studies of free volumes in HPT-deformed Cu, first of all a comparison with recent results on HPT-deformed Ni [9,15,17] is instructive.
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The Ni samples were HPT-deformed under identical conditions as in the present case for Cu, and the sample purity (99.99+wt.%) was also similar.
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As shown in Fig. 6, the dilatometric change in HPT-Ni exhibits a qualitatively similar three-stage behavior.
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In particular, a well-defined narrow stage B also occurs in HPT-Ni due the removal of grain boundaries in the wake of pronounced crystallite growth in this stage [9], similar to Cu.
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However, a number of distinct differences between HPT-Ni and Cu should be noted:
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Stage B in Cu is shifted to lower temperatures by about 40K compared to Ni for similar heating rates.
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(ii)
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The crystallites of HPT-deformed Cu exhibit an isotropic shape, in contrast to HPT-Ni, where a pronounced elongation of the crystallites in the direction tangential to the HPT disk occurs, giving rise to a strong variation in the dilatometric length change with measuring direction, unlike in Cu.
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More than 50% of the total length change in HPT-Cu actually occurs in stage C, whereas for HPT-Ni the regime beyond stage B makes only a minor contribution.
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Stage A, on the other hand, is slightly more pronounced in HPT-Ni ((Δℓ/ℓ0)Stage A=1.57×10-4, mean value of 14 samples) than in Cu ((Δℓ/ℓ0)Stage A=(1.32±0.46)×10-4).
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The shift of the crystal growth-associated stage B towards lower temperatures in Cu compared to Ni (item i; Fig. 6) reflects the different melting temperatures.
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The same is the case for the activation energies of Q=0.99 and 1.20eV determined for Cu and Ni, respectively, from Kissinger analysis (Cu, present studies; Ni [15]) and from Johnson-Mehl-Avrami-Kolmogorov analysis (Ni [15]).
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The value of Q=0.99eV for Cu is identical to that reported by Číček et al. [18] (1.0eV) for HPT-Cu.
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Lower values were reported by Setman et al. [13] (0.48-0.78eV), Cao et al. [19] (0.8eV), and by Molodova et al. [20] (0.68eV), depending on the applied shear strain.
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A remarkably higher value of 1.68eV, caused by impurities of oxygen and phosphorus segregating at grain boundaries, was observed by Amouyal et al. [21].
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In addition to this issue of kinetics, one should point out the major novel aspect in characterizing SPD materials by means of dilatometry: namely, the access to the absolute concentration of free volumes, and in particular to the grain boundary excess volume eGB, the structural key parameter of grain boundaries.
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For Cu, a value (0.46±0.10)×10-10m was deduced from the present measurements, which is slightly higher than recently reported results for Ni [9].
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In the case of Ni with elongated crystallites (item ii), nearly identical values of eGB=0.35×10-10 and 0.32×10-10m were found for the dilatometric measuring directions perpendicular to and parallel to the crystallite elongation, respectively, which confirms the attribution of the length change in stage B to grain boundaries [9].
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As any relaxation processes of the SPD-generated grain boundaries should be finished at the elevated temperatures of stage B, these values of the excess volume can be considered as characteristic values for grain boundaries of polycrystalline Cu and Ni in general.
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It should be noted that the grain boundary excess volume eGB represents the GB expansion with respect to a perfect crystal lattice and should not be intermixed with the grain boundary width δ, which is usually in the range of 0.5nm, i.e. much larger than eGB.
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Only a few experimental data are available in the literature for grain boundary expansion, primarily for isolated grain boundaries with a distinct orientation relation.
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From high-resolution transmission electron microscopy, values for Au of eGB=(0.04-0.10)×10-10m (Ref.
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[22]) or eGB=0.12×10-10m (Ref.
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[23]) are reported.
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From measurements of the grain boundary contact angle in an Al tricrystal, a value eGB=0.64×10-10m is reported by Shvindlerman et al. [24] applying a thermodynamic model.
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For nanocrystalline Pd [25] and Fe [26], values of eGB=0.23×10-10m and 0.19×10-10m were determined from density measurements and modeling of grain growth kinetics, respectively.
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A number of computer simulations of grain boundaries [27-30] deal with the issue of grain boundary expansion.
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Here, however, the choice of the interatomic potentials was found to have a substantial influence on the numerical results [27].
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Most recently, grain boundary expansion data have been reported from molecular dynamics simulations on Ni, eGB=(0.28-0.42)×10-10m for random high-angle grain boundaries [30] and eGB=(0.39-0.41)×10-10m (at T=1200K) for Σ5 grain boundaries [29].
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Here, the matching of the data values with the data of the dilatometric studies of Cu and Ni is remarkable.
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It is worthwhile mentioning a model proposed by Estrin et al. [31], according to which the annealing out of grain boundary excess volume gives rise to linear grain growth instead of a parabolic behavior.
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oastm_01
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From this point of view, it would be interesting to extend the dilatometry to isothermal measurements and to measure the crystallite size in stage B in more detail, in order to derive experimental information on the correlation between the grain growth kinetics and the free volume release.
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oastm_01
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The different dilatometric characteristics of Cu and Ni with respect to stage A and, particularly, stage C (items iii and iv) is considered to arise primarily from the different behaviors of lattice vacancies in both fcc metals.
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oastm_01
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Lattice vacancies generated in highly abundant concentrations by HPT deformation become mobile at about 360K in Ni [32], i.e. in the regime of stage A, whereas in Cu lattice vacancies are already mobile below ambient temperature [33].
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oastm_01
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Therefore, lattice vacancies in Cu may anneal out or form more stable vacancy agglomerates during deformation.
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oastm_01
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oastm_01
This may explain the reduced amplitude of stage A in Cu compared to Ni.
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oastm_01
train5
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oastm_01
For a discussion of the substantial release of free volumes (Δl/l0)Stage C=(3.77±1.33)×10-4 in stage C, one must first note that the removal of remnant grain boundaries in this temperature range contributes to only a minor extent.
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oastm_01
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oastm_01
Taking into account the grain boundary excess free volume derived from stage B, a length change (Δl/l0)Stage C, GB=0.56×10-4 due to the removal of grain boundaries in stage C is derived from the mean crystallite sizes at the onset and the end of this stage, which corresponds to only 15% of the total length change (Δl/l0)Stage C.
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Thus, the major part of stage C is considered to arise from the shrinkage of nanovoids at high temperatures.
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Indeed, the shrinkage of nanovoids in Cu by self-diffusion in this temperature range is well documented from early studies of transmission electron microscopy of coarse-grained Cu in which nanovoids were generated by precipitation of quenched-in vacancies (see Bowden and Balluffi [34]).
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Annealing of vacancy agglomerates in Cu at these elevated temperatures is also deduced from residual resistivity measurements after quenching [33].
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Nanovoids in equal-channel angular pressing (ECAP)-prepared Ti and HPT-prepared Cu were detected by small-angle neutron scattering [35] and positron annihilation [18], respectively.
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Most recently, evidence of percolating porosity in HPT-prepared Cu with a high volume fraction ΔV/V0 in the range of 2-3×10-3 was deduced from radiotracer diffusion and permeation experiments [36].
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Assuming isotropic annealing, the dilatometric length change in stage C minus the contribution from removal of grain boundaries corresponds to a change volume ΔV/V0=3×Δℓ/ℓ0 in the range of 1×10-3, which is smaller than that from the radiotracer experiments by a factor of 2-3.
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The different amounts of porosity estimated from the tracer technique and in the present dilatometric annealing experiments therefore indicate that only part of the porosity is annealed out up to the maximum annealing temperature of 673K - especially since the percolating porosity detected by the tracer method may be considered as the lower limit of the total porosity.
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Indeed, the dilatometric annealing curve shows that the absolute value of the derivative d/dT(ΔV/V0)=3×Δl/l0 is still increasing at the maximum annealing temperature, i.e. the maximum reaction rate of stage C is not yet attained (Fig. 2).
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Dilatometry experiments up to higher temperatures may clarify to what extent porosity can be further removed by annealing.
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The much less pronounced length change in stage C for HPT-deformed Ni (cf.
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Fig. 6) may indicate a substantially reduced amount of porosity or at least a higher thermal stability of such porosity for Ni compared to Cu.
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Tracer diffusion and permeation data for ECAP-prepared metals reveal a higher receptivity for percolating porosity in the case of Cu [37] compared to Ni [8].
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This may be related to the aforementioned different vacancy characteristics, which in the case of Cu may favor the formation of stable vacancy agglomerates during deformation.
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In conclusion, the direct and specific method of high-precision dilatometry has proven to be a powerful tool for the study of free volume in bulk nanocrystalline metals.
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In addition to issues of defect kinetics, the absolute value of free volumes, such as the grain boundary excess volume, can be measured directly.
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Acknowledgement
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appreciated.
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Preparation and characterization of a dual-layer carbon film on 6H-SiC wafer using carbide-derived carbon process with subsequent chemical vapor deposition
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Abstract
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It is reported that a dual-layer carbon film on SiC wafer is prepared using carbide-derived carbon (CDC) process with subsequent chemical vapor deposition (CVD).
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oastm_02
train4
semi
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The dual-layer film includes a sub-layer of CDC and a top layer of CVD, which are prepared by chlorination of SiC and pyrolysis of CCl4 at high temperature respectively.
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The CDC and CVD layers are mainly amorphous.
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oastm_02
train4
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oastm_02
And similar dispersion effects are observed in the Raman spectra, although the D-band position of the CVD layer shifts to higher wavenumber (~1354cm-1) than that of the CDC layer (~1337cm-1).
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oastm_02
train4
semi
oastm_02
Surface chemistry analysis suggests that the unstable chemical bonds, mainly C-Cl, as well as dangling bonds in the CDC layer play an important role in promoting the nucleation of CVD carbon.
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The surface morphology evolvement from SiC wafer to CDC layer and to dual-layer film is investigated by atomic force microscopy [AFM] and field emission scanning electronic microscopy [FESEM].
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oastm_02
train4
semi
oastm_02
The nanoporous surface formed in the CDC process is favorable for capturing carbon species from the gas phase and can act as a "seedbed" for the nucleation and growth of CVD layer.
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oastm_02
train4
semi
oastm_02
The primary tribological study indicates that the dual-layer film shows great advantages in friction reduction and wear resistance with comparison to SiC and CDC layer, suggesting its potential in lubrication for SiC-based moving components.
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