(Zahid Hussain)
1
(Hye-Won Yang)
1
(Byung-Sang Choi)
12*
Copyright © 2021 The Korean Institute of Metals and Materials
Key words(Korean)
hexagonal boron nitride, reinforcement, 3Di-hBN-Cu-Ni composite, MOCVD
1. Introduction
The development of two-dimensional (2D) materials has opened up the possibilities
for their application to improve the properties of metals and alloys [1-8], because 2D materials have the ability to alter the properties of metals at the
nanoscale. A single layer of hexagonal boron nitride (hBN) is structurally similar
to graphene (carbon system) with the hexagonal lattices occupied by boron and nitrogen
atoms. hBN has a lattice parameter of 25 nm and possesses extraordinary properties,
including high chemical stability [9], high mechanical strength [9], low density [10], high thermal stability [11], and high thermal shock resistance [12]. These excellent properties can be utilized to improve the performance of various
metal matrix composites (MMCs) by constructing three-dimensionally interconnected
(3Di) hBN layers in their grain boundaries. Similar approaches have been employed
by other researchers, who have used 3D-networked graphene to tailor the properties
of MMCs [8,13-15]. For instance, Chen et al. [13] enhanced the yield and tensile strengths of copper by wrapping graphene around copper
grains using chemical vapor deposition (CVD). They reported that the graphene acted
as a barrier to dislocation movement, and consequently, the elastic modulus and strength
were improved. Li et al. [8] reported that Cu-graphene composites had a higher thermal conductivity than pure
copper because graphene offered an effective path for heat transfer between the Cu
grain boundaries. Other properties, such as corrosion resistance and wear resistance,
have also been improved [16,17].
Because of its structure similarity to graphene, introducing the hBN into metal matrices
can result in similar effects. Several researchers have used boron nitride nanoparticles
to enhance the strength, hardness, wear, and corrosion resistance of metallic alloys
[18-22]. For instance, the microstructure and properties of BN/Ni-Cu composites fabricated
by powder technology were reported by Tantaway et al. [23]. They found that the BN content led to a decrease in density and an increase in
the hardness, electrical resistivity, and saturation magnetization of the composite.
Omayma et al. [10] fabricated Cu/hBN nanocomposites by the PM route, in which powder mixtures of Cu
and hBN were compacted and sintered at various temperatures ranging from 950 °C to
1000 °C. They found that the physical, mechanical and tribological properties of the
composite were influenced by the hBN. However, we have not found any published studies
in which metal matrices were reinforced by a 3Di network of hBN layers.
Recently, various techniques have been utilized to fabricate reinforced MMCs by incorporating
graphene. For instance, Xiong et al. [24] introduced graphene in Cu by the reduction of reduced graphene oxide through sintering.
Similarly, ball milling, molecular-level synthesis, spark plasma sintering, and epitaxial
growth have been used to improve the strength of composites, employing graphene as
a reinforcement [13, 25-28]. However, each of the these strengthening techniques has some limitations. For instance,
ball milling and molecular-level mixing may allow a uniform dispersion of the reinforcement
material but may impart structural defects due to shear stress and contamination during
the fabrication process [29]. A well-ordered/-aligned, uniformly dispersed, and continuous graphene network is
essential to attain the best reinforcement results [13]. Kawk et al. [30] introduced a simple, economically efficient two-step process with the potential
to deliver better-quality products with uniformly dispersed and continuous graphene
networks.
The two-step process involves the compaction of a metallic powder followed by CVD.
In this study, we fabricated a 3Di-hBN-Cu-Ni composite using a similar simple two-step
process. Various characterization techniques were employed to confirm the formation
of 3Di-hBN surrounding the grains of the Cu-Ni alloys. Cu-Ni-based alloys have been
employed in various industries, including shipbuilding, construction, and processing,
because of their high mechanical strength and corrosion resistance at elevated temperatures.
The 3Di-hBN-Cu-Ni composite is expected to deliver better corrosion, mechanical, and
wear characteristics than the Cu-Ni alloy. Moreover, the 3Di-hBN, which has a foam-like
3D porous structure, was separated from the 3Di-hBN-Cu-Ni composite and could have
further applications in the fields of biomedicine, electronics, and energy storage
[31-33].
2. Experimental Procedures
2.1 Fabrication of 3Di-hBN-Cu-Ni Composite
Cu powder (99.5% purity) with spheroidal particles of size 14-25 µm and Ni powder
(>99.5% purity) with spheroidal particles of size ~1 µm were purchased from Sigma-Aldrich
and used after heat treatment (200 °C for 2 h in an H2 environment) to remove any moisture or oxide contents. The chemical compositions
of Cu and Ni powders are provided in Table 1. Cu and Ni powders (70 wt.% Cu, 30 wt.% Ni) were mixed manually using mortar, while
taking care not to change the particle size distribution. The mixture was compacted
in a mold using a double-action oil hydraulic press at compaction pressures of 60,
110, 220, 280, 335, and 390 MPa. The exertion of high pressure on the spheroidal particles
caused mechanical cold locking among the particles, forming a compact disc with the
approximate diameter and thickness of 15 mm and 1.2 mm, respectively. As shown on
the fracture surface of the cross section of the compact disc in Figure 1, the relatively large Cu particles induced mechanical interlocking, due to their
deformation, and Ni particles filled the gaps between the Cu particles. The discs
were then placed in a quartz glass tube furnace with a tube diameter of 23 mm for
metal-organic CVD (MOCVD). The compaction pressure and sintering time were varied
to determine the optimum conditions for the fabrication of 3D-hBN in the 3Di-hBN-Cu-Ni
composite.
Figure 2 (a) shows a schematic of the system used to fabricate the 3Di-hBN-Cu-Ni composites. First,
the system was flushed with argon at least three times to remove air from the MOCVD
tube. The furnace temperature was increased to 400 °C at a rate of 16.6 °C/min and
then maintained constant in a hydrogen environment at 330 Torr for 1 h for deoxidation.
Subsequently, the temperature was raised to 1000 °C at the same rate and maintained
constant for 15 or 30 min. Finally, MOCVD was performed for 15 min at 450 Torr using
heated decaborane (B10H14) as the boron source and ammonia (NH3) as the nitrogen source. Decaborane was the preferred boron precursor because of
its (i) easy handling, (ii) commercial availability, and (iii) stability, which minimized
the formation of undesired side products at elevated temperatures that could potentially
decrease hBN yield [34]. Decaborane is a crystalline solid with a melting temperature of 98-100 °C and its
vapor pressure can be easily controlled by varying the temperature from room temperature
to 100 °C. At approximately 100 °C, the vapors produced upon evaporation can be transported
into the MOCVD growth zone by an inert carrier gas (Ar) at a flow rate of 1 sccm.
Ammonia gas was introduced as a nitrogen source in the MOCVD reaction zone at a flow
rate of 2 sccm. At 1000 °C, ammonia and decaborane dissociated into nitrogen and boron
atoms, respectively.
The entire process (heating, sintering, and MOCVD) was conducted at a hydrogen flow
rate of 10 sccm. The 3Di-hBN-Cu-Ni composite, fabricated using the simple two-step
process, is shown in Figure 2(b).
2.2 Characterization
Optical microscopy (OM) and scanning electron microscopy (SEM) investigations of the
3Di-hBN-Cu-Ni composite samples were conducted after the sample was cut in half and
the cutting surface was polished with emery papers down to 4000 grit. Finally, the
polished surface was etched at room temperature using a mixed solution of 1 M FeCl3 and 0.1 M HCl to reveal the microstructure [30]. For transmission electron microscopy (TEM) investigations, the 3Di-hBN-Cu-Ni composite
samples were mechanically polished to a thickness of 100 µm and cut into small pieces
of 3 mm diameter. Then, the Cu-Ni was etched out, leaving only 3Di-hBN foam, which
was transferred after thorough cleaning to the TEM grid for investigation. A qualitative
phase analysis of the 3Di-hBN-Cu-Ni composite was performed by XRD analysis using
Cu Kα radiation with a wavelength of 1.54 Å and a scanning angle of 20 °-100 °. The density
of the 3Di-hBN-Cu-Ni composite samples was measured using the Archimedes immersion
technique.
2.3 Three-Dimensionally Interconnected hBN
The 3Di-hBN-Cu-Ni composite samples were cut into small pieces, polished, and placed
in an etchant for a sufficient duration to etch out Cu-Ni completely, so that only
3Di-hBN remained. Then, the foam-like 3Di-hBN samples were removed and washed several
times with deionized water. To obtain a stable 3D structure of the 3Di-hBN, the freeze-drying
method was used to ensure that there was no effect of liquid capillary force, and
that the 3Di-hBN did not structurally collapse [35].
3. Results and Discussions
Figure 3 shows a schematic of the processes involved in the synthesis of the 3Di-hBN-Cu-Ni
composite. During the sintering, a reduction in volume and the formation of a Cu-Ni
solid solution occur due to the diffusion of metals under the driving force, to reduce
the excess surface energy [36,37]. Consequently, the overall volume of the compact disc is reduced and densification
occurs. The formation of 3Di-hBN in the composite is likely to occur in three stages
[8]. First, the diffusion of metal occurs to reduce the surface energy, resulting in
the formation of large particles (consolidation). At the same time, the diffusion
of Ni to Cu or vice versa occurs to form a solid solution of Cu-Ni. Next, during the
MOCVD process, the dissociation of ammonia and decaborane produces nitrogen and boron
atoms that diffuse into the Cu-Ni alloy at 1000 °C. Finally, upon cooling, the nitrogen
and boron atoms precipitate out and alternately join together to form 2D hBN layer(s)
along the interfaces of the Cu-Ni grains [38] resulting in the formation of the Cu-Ni composite.
Small pores or voids can form during the sintering, as Cu and Ni particles grow to
reduce their surface energy. This is probably due to insufficient sintering time or
excessive free space among the particles. These pores may also act as catalytic sites
for the nucleation and growth of bulk hBN. The small lighter grey areas (indicated
by small white loops) in Figure 4 indicate the bulk hBN that accumulated on the pores during the MOCVD process. These
pores, generated during the sintering process and then filled with bulk hBN during
the MOCVD process, are undesirable, as they may adversely affect the mechanical, thermal,
and wear characteristics of the composite. Therefore, the processing parameters, such
as compaction pressure and sintering time, must be varied to determine the optimal
conditions for fabricating 3Di-hBN-Cu-Ni composites without the formation of bulk
hBN.
Figure 5 shows the density of the 3Di-hBN-Cu-Ni composite as a function of the compaction
pressure and sintering time. The density of the composite increased with increasing
compaction pressure. However, beyond a certain compaction pressure, the density decreased.
This trend occurred because at pressures below 280 MPa, the compaction pressure was
not enough, resulting in a low density of compaction. This consequently resulted in
a lower density composite with voids or pores. Although these pores were filled with
bulk hBN during the subsequent MOCVD process, the overall density of the composite
could not be increased because the density of hBN (2.1 g/cm3) is significantly lower than that of Cu-Ni (~8.9 g/cm3). On the other hand, during compaction at high pressures (>280 MPa), the particles
on the surface were pressed with a relatively greater force than those inside the
compact disc, because of friction between the particles. Consequently, at pressures
exceeding 280 MPa, the surface particles of the compact disc were denser than the
inner particles. The inner particles, having a longer diffusion distance owing to
their lower density, resulted in the formation of pores due to insufficient diffusion
or short sintering time. This led to the relatively larger size of pores, or shrinkage,
as indicated in the OM images in Figure 6(b). Hence, the density of the 3Di-hBN-Cu-Ni composite was slightly lower at higher pressures
(>280 MPa), as shown in Figure 5. This is also evident in the OM images shown in Figure 4. The white arrows in Figure 4 indicate the bulk hBN present in the microstructure of the 3Di-hBN-Cu-Ni composite.
A relatively higher volume fraction of bulk hBN was observed when the compaction pressure
was lower or higher than 280 MPa. Furthermore, the density of the 3Di-hBN-Cu-Ni composite
also depends on the sintering time, as shown in Figure 5. A longer sintering time led to fewer pores (i.e., a lower volume fraction of bulk
hBN), and consequently more densification occurred.
The 3Di-hBN-Cu-Ni-hBN composite fabricated under the optimized conditions (a compaction
pressure of 280 MPa and a sintering time of 30 min) was examined using SEM. While
some of the bulk hBN was removed during the polishing and etching, the SEM image in
Figure 7 and the EDS results in Table 2 show the Cu-Ni grains, bulk hBN of less than 5 µm in size, and hBN along the interfaces.
The Cu-Ni grains, grain boundaries, and bulk hBN in Figure 7 were analyzed using energy-dispersive X-ray spectroscopy (EDS). Boron and nitrogen
were observed (locations (a) and (b) in Figure 7) in excess along with minute amounts of other impurities, such as silicon, carbon,
and oxygen, as listed in Table 2. These impurities probably entered the structure during the polishing and etching
processes. Location (a) in Figure 7 is a pore that was first formed as a consequence of sintering and then filled with
bulk hBN during the subsequent MOCVD process. Considering the average size (5 µm)
of these sites (location (a)), the presence of bulk hBN was verified through EDS analysis.
Further EDS analysis at the grain boundaries (location (b)) revealed that the grain
boundaries were also mostly occupied by boron and nitrogen with an approximate stochiometric
ratio of 1:1. As expected, the Cu-Ni grains (location (c) in Figure 7) comprised Cu and Ni atoms with a ratio of approximately 7 to 3, as shown in Table 2. The solubility of B and N in Cu-Ni alloy at 1000 °C is very small (~ ppm) [23] and most of the atoms (B and N) precipitated out during cooling, thus forming hBN
with B and N having a stochiometric ratio of 1:1 at the grain boundaries [39-41].
The SEM image in Figure 8 shows various hBN layers that interconnect to form a foam-like structure with pockets
and channels. The channels are the connected areas between the Cu-Ni grains formed
by etching. The average pocket size (10-20 µm) in 3Di-hBN (Figure 8) is approximately equal to the average grain size of the 3Di-hBN-Cu-Ni composite
(Figures 4 and 7) indicating that the hBN layers are wrapped around the Cu-Ni grains in the 3Di-hBN-Cu-Ni
composite.
The XRD pattern of the 3Di-hBN-Cu-Ni composite is shown in Figure 9(a). This pattern only shows the crystalline phase of Cu-Ni solid solution. The elemental
distribution map of the 3Di-hBN-Cu-Ni composite shown in Figure 9(b) shows a uniform distribution of Cu and Ni, indicating the formation of a Cu-Ni solid
solution.
The 3Di-hBN foam was inspected using TEM. The low-magnification bright-field TEM image
in Figure 10(a) shows a complex morphology with curvatures and overlapped structures, where the 3D
layers of hBN (shown in Figure 8) collapsed after their transfer to the TEM grid under the capillary force during
the drying process. The selected-area electron diffraction pattern of 3Di-hBN (inset
in Fig 10(a)) indicates multiple orientations associated with a couple of layers with different
orientations. The high-resolution TEM (HR-TEM) image (Figure10 (b)) reveals 2-6 layers of hBN, with an interlayer distance of approximately 0.25 nm
(inset in Figure 10(b)), which is considered to be the thickness of a single layer of 2D hBN.
4. Conclusion
3Di-hBN-Cu-Ni composites were synthesized via a simple two-step process: (1) compacting
Cu and Ni powder mixtures without any additives and (2) MOCVD. The density of the
composite was the highest (7.75 g/cm3) when the compaction pressure and sintering time were 280 MPa and 30 min, respectively.
OM, SEM, and TEM images indicated that these conditions were optimal for the growth
of interconnected network of hBN in the 3Di-hBN-Cu-Ni composite. SEM investigations
and EDS analysis revealed that the grain boundaries were mostly occupied by boron
and nitrogen atoms. 3Di-hBN was obtained after etching the Cu-Ni alloy. The average
pocket size of the foam was 10-20 µm. The 3Di-hBN-Cu-Ni composite with a density of
7.75 g/cm3 was shown to have a three-dimensional network of hBN. The structural investigation
of 3Di-hBN using TEM revealed 2-6 layers with an interlayer distance of 0.25 nm. This
study can be extended further, with characterizations of the physical and chemical
properties of the 3Di-hBN-Cu-Ni composite and 3Di-hBN.
Acknowledgements
This study was supported by research funds from Chosun University (2020).
REFERENCES
Bartolucci S. F., Paras J., Rafiee M. A., Rafiee J., Lee S., Kapoor D., Koratkar N.,
Mater. Sci. Eng. A,528, 7933 (2011)
Chu K., Jia C., Phys. Status Solidi A,211, 184 (2014)
Guo S., Zhang X., Shi C., Liu E., He C., He F., Zhao N., Mater. Sci. Eng. A,766, 138365
(2019)
Liu Y., Wu H., Chen G., Polym. Composite,37, 1190 (2016)
Rafiee M. A., Rafiee J., Wang Z., Song H., Yu Z. Z., Koratkar N., ACS Nano,3, 3884
(2009)
Yang W., Zhao Q., Xin L., Qiao J., Zou J., Shao P., Yu Z., Zhang Q., Wu G., J. Alloy.
Compd,732, 748 (2018)
Jo Y., Li X., Jo J. M., Choi B. S., J. Adv. Eng. and Tech,12, 91 (2019)
Li X., Ring T. A., Choi B. S., Korean J. Met. Mater,57, 529 (2019)
Singh S. P., MSc Thesis,54-55, Thapar Insitute of Technology (2018)
Omayma A. M., Oqail A. A., Emad M. M., Sheikh M. E., J. Alloy. Compd,625, 309 (2015)
Liu Z., Gong Y., Zhou W., Ma L., Yu J., Idrobo J. C., Jung J., MacDonald A. H., Vajtai
R., Lou J., Ajayan P. M., Nat. commun,4, 1 (2013)
Duan X., Yang Z., Chen L., Tian Z., Cai D., Wang Y., Jia D., Zhou Y., J. Eur. Ceram.
Soc,36, 3725 (2016)
Chen Y., Zhang X., Liu E., He C., Han Y., Li Q., Nash P., Zhao N., J. Alloy. Compd,688,
69 (2016)
Kim S. E., Shon I. J., Korean J. Met. Mater,56, 854 (2018)
Song M. Y., Kwak Y. J., Choi E., Korean J. Met. Mater,56, 524 (2018)
Wu M., Hou B., Shu S., Li A., Geng Q., Li H., Shi Y., Yang M., Du S., Wang J. Q.,
Liao S., Jiang N., Dai D., Lin C. T., Nanomaterials,9, 498 (2019)
Tripathi K., Gyawali G., Lee S. W., ACS Appl. Mater. Inter,9, 32336 (2017)
Gopinath S., Prince M., Raghav G., Mater. Res. Express,7, 016582 (2020)
Reddy A. C., Proc. 5th International Conference on Modern Materials and Manufacturing,
409MMM (2013)
Zitoun E., Reddy A. C., Proc. 6th National Conference on Materials and Manufacturing
Processes, Hyderabad (2008)
Khatavakar A., Mandave A. K., Baviskar D. D., Shinde S. L., Int. Res. J. Eng. Technol,5,
3791 (2018)
Cho W. J., Shon I. J., Korean J. Met. Mater,56, 658 (2018)
El-Tantawy A., Daoush W. M., El-Nikhaily A. E., J. Exp. Nanosci,13, 174 (2018)
Li J. L., Xiong Y. C., Wang X. D., Yan S. J., Yang C., He W.W., Chen J. Z., Wang S.
Q., Zhang X. Y., Dai S. L., Mater. Sci. Eng. A,626, 400 (2015)
Cao M., Xiong D. B., Tan Z., Ji G., Ahmadi B. A., Guo Q., Fan G., Guo C., Li Z., Zhang
D., Carbon,117, 65 (2017)
Jiang R., Zhou X., Fang Q., Liu Z., Mater. Sci. Eng. A,654, 124 (2016)
Wang M., Wang L. D., Sheng J., Yang Z. Y., Shi Z. D., Zhu Y. P., Li J., Fei W. D.,
J. Alloy. Compd,798, 403 (2019)
Yang Z., Wang L., Sji Z., Wang M., Cui Y., Wei B., Xu S., Zhu Y., Fei W., Carbon,127,
329 (2018)
Naseer A., Ahmad F., Aslam M., Guan B. H., S W., Harun W., Muhammad N., Raza M. R.,
German R. M., Mater. Manuf. Process,34, 957 (2019)
Kawk S. R., Ring T. A., Choi B. S., J. Ind. Eng. Chem,70, 484 (2019)
Gautam C., Chakravarty D., Gautam A., Tiwary C. S., Woellner C. F., Mishra V. J. K.,
Ahmad N., Ozden S., Jose S., Biradar S., Vajtai R., Trivedi R., Galvao D. S., Ajayan
P. M., ACS Omega,3, 6013 (2018)
Guiney L. M., Mansukhani N. D., Jakus A. E., Wallace S. G., Shah R. N., Hersam M.
C., Nano Lett,18, 3488 (2018)
Yin J., Li X., Zhou J., Guo W., Nano Lett,13, 3232 (2013)
Chatterjee S., Kim M. J., Zakhrov D. N., Kim S. M., Stach E. A., Maruyama B., Sneddon
L. G., Chem. Mater,24, 2872 (2012)
Fang Q., Shen Y., Chen B., Chem. Eng. J,264, 753 (2015)
German R. M., Crit. Rev. Solid State Mater. Sci,35, 263 (2010)
McDonald S. A., Holzner C., Lauridsen E. M., Reischig P., Merkle A. P., Withers P.
J., Sci. Rep,7, 1 (2017)
Joshi S., Ecija D., Koitz R., Iammuzzi M., Seitsonen A. P., Hutter J., Sachdev H.,
Vijayaraghavan S., Bischoff F., Seufert K., Barth J. V., Auwarter W., Nano lett,12,
5821 (2012)
Koepke J. C., Wood J. D., Chen Y., Schucker S. W., Liu X., Chang N. N., Niwnhaus L.,
Do J. W., Carrion E. A., Hewaparakrama J., Rangarajan A., Datye I., Mehta R., Haasch
R. T., Gruebele M., Girolami G. S., Pop E., Lyding J. W., Chem. Mater,28, 4169 (2016)
Chakrabarty K., Arnold I., Catledge S. A., J. Vac. Sci. Technol. A,37, 061507-1 (2019)
Khan A. F., Randviir E. P., Brownson D. A. C., Ji X., Smith G. C., Banks C. E., Electroanal,29,
622 (2017)
Figures and Tables
Fig. 1.
Cu and Ni particles shown at the fracture surface of the disc after compaction.
Fig. 2.
(a) Schematic illustration of the fabrication of 3Di-hBN-Cu-Ni composite and (b) disc-shaped
3Di-hBN-Cu-Ni composite fabricated using a simple two-step process.
Fig. 3.
Schematic showing the process of formation of 3Di-hBN-Cu-Ni composite.
Fig. 4.
OM images of 3Di-hBN-Cu-Ni composites under various conditions.
Fig. 5.
Density of the 3Di-hBN-Cu-Ni composite as a function of compacting pressure and sintering
time.
Fig. 6.
OM images of 3Di hBN-Cu-Ni composites processed at compaction pressures of (a) 280
MPa and (b) 335 MPa.
Fig. 7.
SEM image showing the surface morphology of the 3Di-hBN-Cu-Ni composite fabricated
with compaction pressure of 280 MPa and sintering time of 30 min.
Fig. 8.
SEM image showing the 3D interconnected network of hBN in 3Di hBN-Cu-Ni composite
produced under compaction pressure of 280 MPa and sintering time of 30 min, respectively.
Fig. 9.
(a) XRD pattern of 3Di-hBN-Cu-Ni composite and (b) Elemental distribution map of 3Di-hBN-Cu-Ni
composite.
Fig. 10.
TEM investigation: (a) low-magnification bright-field TEM image of 3Di-hBN and (b)
HR-TEM image showing 2–6 layers of hBN (inset shows a distance of 0.25 nm between
layers).
Table 1.
Chemical compositions of Cu and Ni powders.
|
Material
|
Purity
|
Trace metals in ppm
|
|
Cu Powder
|
> 99.5%
|
Fe 80.0, Na 9.42, Mn 7.6, Mg 4.69, Al 4.4, and B 1.98
|
|
Ni Powder
|
> 99.5%
|
Ag 1.3, Al17.9, Ba 0.8 Ca 27.9, Cr 3.6, Cu 305.2, Fe 383.7 Mg 2.0, Mn 2.6, Na 11.1,
Pd 8.0, Ti 144.5, and V 25.4
|
Table 2.
EDS results of the 3Di-hBN-Cu-Ni composite for the microstructure shown in Figure
7.
|
Element
|
Location
|
Cu
|
Ni
|
B
|
N
|
Si
|
C
|
O
|
|
At. %
|
(a) in Figure 7
|
-
|
-
|
49.32
|
47.45
|
0.76
|
0.87
|
0.68
|
|
At. %
|
(b) in Figure 7
|
20.51
|
8.32
|
34.26
|
34.82
|
0.89
|
0.57
|
0.78
|
|
At. %
|
(c) in Figure 7
|
71.28
|
28.42
|
-
|
-
|
-
|
-
|
-
|