Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites
The nanoscale mono- (powders) and complex (compacted nanocomposites) carbides of d-transition metals are synthesized by mechanical alloying in a high-energy planetary ball mill from a charge containing the carbon nanotubes. The effect of multiwalled carbon nanotubes on reaction milling of the obtain...
Saved in:
| Published in: | Успехи физики металлов |
|---|---|
| Date: | 2019 |
| Main Authors: | , , , , |
| Format: | Article |
| Language: | English |
| Published: |
Інститут металофізики ім. Г.В. Курдюмова НАН України
2019
|
| Online Access: | https://nasplib.isofts.kiev.ua/handle/123456789/167921 |
| Tags: |
Add Tag
No Tags, Be the first to tag this record!
|
| Journal Title: | Digital Library of Periodicals of National Academy of Sciences of Ukraine |
| Cite this: | Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites / O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, V.V. Zavodyannyi, N.N. Belyavina // Progress in Physics of Metals. — 2019. — Vol. 20, No 1. — P. 5-51. — Bibliog.: 65 titles. — eng. |
Institution
Digital Library of Periodicals of National Academy of Sciences of Ukraine| id |
nasplib_isofts_kiev_ua-123456789-167921 |
|---|---|
| record_format |
dspace |
| spelling |
Nakonechna, O.I. Dashevskyi, M.M. Boshko, О.І. Zavodyannyi, V.V. Belyavina, N.N. 2020-04-16T15:14:44Z 2020-04-16T15:14:44Z 2019 Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites / O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, V.V. Zavodyannyi, N.N. Belyavina // Progress in Physics of Metals. — 2019. — Vol. 20, No 1. — P. 5-51. — Bibliog.: 65 titles. — eng. 1608-1021 DOI: https://doi.org/10.15407/ufm.20.01.005 https://nasplib.isofts.kiev.ua/handle/123456789/167921 The nanoscale mono- (powders) and complex (compacted nanocomposites) carbides of d-transition metals are synthesized by mechanical alloying in a high-energy planetary ball mill from a charge containing the carbon nanotubes. The effect of multiwalled carbon nanotubes on reaction milling of the obtained materials is analyzed. The features of formation mechanism of metal carbides at the mechanical alloying are clarified. Механохімічним методом у високоенергетичному планетарному кульовому млині з шихти, що містить вуглецеві нанотрубки, синтезовано нанорозмірні моно-(порошки) та подвійні (компактовані нанокомпозити) карбіди d-перехідних металів. Розглянуто вплив багатошарових вуглецевих нанотрубок на механохімічний синтез одержаних матеріалів. З’ясовано особливості механізму формування карбідів перехідних металів у процесі механохімічного синтезу. Механохимическим методом в высокоэнергетической планетарной шаровой мельнице из шихты, содержащей углеродные нанотрубки, синтезированы наноразмерные моно- (порошки) и двойные (компактированные нанокомпозиты) карбиды d-переходных металлов. Рассмотрено влияние многослойных углеродных нанотрубок на механохимический синтез полученных материалов. Выяснены особенности механизма формирования карбидов переходных металлов в процессе механохимического синтеза. The authors appreciate sincerely Prof. V. Tkach and Dr. D. Stratiichuk (V. N. Bakul Institute for Superhard Material of the National Academy of Sciences of Ukraine), L. Kapitanchuk (Paton Electric Welding Institute of the National Academy of Science of Ukraine) and Prof. M. Semen’ko (Department of Physics, Taras Shevchenko National University of Kyiv) for their help in the preparation of this manuscript and fruitful discussions. en Інститут металофізики ім. Г.В. Курдюмова НАН України Успехи физики металлов Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites Вплив вуглецевих нанотрубок на механохімічний синтез нанопорошків карбідів d-металів і нанокомпозитів на їх основі Влияние углеродных нанотрубок на механохимический синтез нанопорошков карбидов d-металлов и нанокомпозитов на их основе Article published earlier |
| institution |
Digital Library of Periodicals of National Academy of Sciences of Ukraine |
| collection |
DSpace DC |
| title |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| spellingShingle |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites Nakonechna, O.I. Dashevskyi, M.M. Boshko, О.І. Zavodyannyi, V.V. Belyavina, N.N. |
| title_short |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| title_full |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| title_fullStr |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| title_full_unstemmed |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| title_sort |
effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites |
| author |
Nakonechna, O.I. Dashevskyi, M.M. Boshko, О.І. Zavodyannyi, V.V. Belyavina, N.N. |
| author_facet |
Nakonechna, O.I. Dashevskyi, M.M. Boshko, О.І. Zavodyannyi, V.V. Belyavina, N.N. |
| publishDate |
2019 |
| language |
English |
| container_title |
Успехи физики металлов |
| publisher |
Інститут металофізики ім. Г.В. Курдюмова НАН України |
| format |
Article |
| title_alt |
Вплив вуглецевих нанотрубок на механохімічний синтез нанопорошків карбідів d-металів і нанокомпозитів на їх основі Влияние углеродных нанотрубок на механохимический синтез нанопорошков карбидов d-металлов и нанокомпозитов на их основе |
| description |
The nanoscale mono- (powders) and complex (compacted nanocomposites) carbides of d-transition metals are synthesized by mechanical alloying in a high-energy planetary ball mill from a charge containing the carbon nanotubes. The effect of multiwalled carbon nanotubes on reaction milling of the obtained materials is analyzed. The features of formation mechanism of metal carbides at the mechanical alloying are clarified.
Механохімічним методом у високоенергетичному планетарному кульовому млині з шихти, що містить вуглецеві нанотрубки, синтезовано нанорозмірні моно-(порошки) та подвійні (компактовані нанокомпозити) карбіди d-перехідних металів. Розглянуто вплив багатошарових вуглецевих нанотрубок на механохімічний синтез одержаних матеріалів. З’ясовано особливості механізму формування карбідів перехідних металів у процесі механохімічного синтезу.
Механохимическим методом в высокоэнергетической планетарной шаровой мельнице из шихты, содержащей углеродные нанотрубки, синтезированы наноразмерные моно- (порошки) и двойные (компактированные нанокомпозиты) карбиды d-переходных металлов. Рассмотрено влияние многослойных углеродных нанотрубок на механохимический синтез полученных материалов. Выяснены особенности механизма формирования карбидов переходных металлов в процессе механохимического синтеза.
|
| issn |
1608-1021 |
| url |
https://nasplib.isofts.kiev.ua/handle/123456789/167921 |
| citation_txt |
Effect of carbon nanotubes on mechanochemical synthesis of d-metal carbide nanopowders and nanocomposites / O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, V.V. Zavodyannyi, N.N. Belyavina // Progress in Physics of Metals. — 2019. — Vol. 20, No 1. — P. 5-51. — Bibliog.: 65 titles. — eng. |
| work_keys_str_mv |
AT nakonechnaoi effectofcarbonnanotubesonmechanochemicalsynthesisofdmetalcarbidenanopowdersandnanocomposites AT dashevskyimm effectofcarbonnanotubesonmechanochemicalsynthesisofdmetalcarbidenanopowdersandnanocomposites AT boshkooí effectofcarbonnanotubesonmechanochemicalsynthesisofdmetalcarbidenanopowdersandnanocomposites AT zavodyannyivv effectofcarbonnanotubesonmechanochemicalsynthesisofdmetalcarbidenanopowdersandnanocomposites AT belyavinann effectofcarbonnanotubesonmechanochemicalsynthesisofdmetalcarbidenanopowdersandnanocomposites AT nakonechnaoi vplivvuglecevihnanotruboknamehanohímíčniisinteznanoporoškívkarbídívdmetalívínanokompozitívnaíhosnoví AT dashevskyimm vplivvuglecevihnanotruboknamehanohímíčniisinteznanoporoškívkarbídívdmetalívínanokompozitívnaíhosnoví AT boshkooí vplivvuglecevihnanotruboknamehanohímíčniisinteznanoporoškívkarbídívdmetalívínanokompozitívnaíhosnoví AT zavodyannyivv vplivvuglecevihnanotruboknamehanohímíčniisinteznanoporoškívkarbídívdmetalívínanokompozitívnaíhosnoví AT belyavinann vplivvuglecevihnanotruboknamehanohímíčniisinteznanoporoškívkarbídívdmetalívínanokompozitívnaíhosnoví AT nakonechnaoi vliânieuglerodnyhnanotruboknamehanohimičeskiisinteznanoporoškovkarbidovdmetallovinanokompozitovnaihosnove AT dashevskyimm vliânieuglerodnyhnanotruboknamehanohimičeskiisinteznanoporoškovkarbidovdmetallovinanokompozitovnaihosnove AT boshkooí vliânieuglerodnyhnanotruboknamehanohimičeskiisinteznanoporoškovkarbidovdmetallovinanokompozitovnaihosnove AT zavodyannyivv vliânieuglerodnyhnanotruboknamehanohimičeskiisinteznanoporoškovkarbidovdmetallovinanokompozitovnaihosnove AT belyavinann vliânieuglerodnyhnanotruboknamehanohimičeskiisinteznanoporoškovkarbidovdmetallovinanokompozitovnaihosnove |
| first_indexed |
2025-11-24T20:50:57Z |
| last_indexed |
2025-11-24T20:50:57Z |
| _version_ |
1850493158060720128 |
| fulltext |
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20 No. 1 5
© o.I. NAKoNechNA, M.M. DASheVSKyI, О.І. BoShKo,
V.V. ZAVoDyANNyI, N.N. BelyAVINA, 2019
https://doi.org/10.15407/ufm.20.01.005
O.I. NakONechNa 1, M.M. DashevskyI 1, О.І. BOshkO 2,
v.v. ZavODyaNNyI 3, and N.N. BelyavINa1
1 Taras Shevchenko National University of Kyiv,
4 Glushkov Ave., UA-03022 Kyiv, Ukraine
2 G.V. Kurdyumov Institute for Metal Physics of the N.A.S. of Ukraine,
36 Academician Vernadsky Blvd, UA-03142 Kyiv, Ukraine
3 Kherson State Agrarian University,
23 Stritenska Str., UA-73006 Kherson, Ukraine
effect of carbon nanotUbes
on MechanocheMical synthesis
of d-Metal carbide nanoPowders
and nanocoMPosites
the nanoscale mono- (powders) and complex (compacted nanocomposites) carbides of
d-transition metals are synthesized by mechanical alloying in a high-energy planetary
ball mill from a charge containing the carbon nanotubes. the effect of multiwalled
carbon nanotubes on reaction milling of the obtained materials is analyzed. the
features of formation mechanism of metal carbides at the mechanical alloying are
clarified. Particularly, as is shown, at the first stage of the synthesis (up to 60 min
of processing of the charge in a ball mill), the amorphization of the carbon nanotubes
and crushing of particles of the initial metal along the grain boundaries occurs
concurrently. then, the amorphous carbon enters into the metal lattice forming an
interstitial solid solution, resulting in deformation of the metal crystal lattice. At
the second stage of synthesis (from 60 to 250 minutes of processing), the process of
embedding of the carbon atoms in metal matrix is accelerated and the formation of
the carbide phases on surface of the parent metal particles begins. the third stage
of synthesis completes the formation of carbide. As revealed, the processing time
required for the complete transformation of the initial components to the carbide
correlates with the enthalpy of its formation, and the fields of mechanical stress are
relaxed over two main channels: heating and grinding. As found out, the carbides
of d-transition metals studied in this work are formed mainly due to self-supporting
6 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
1. Introduction: Mechanochemical Processing
1.1. Brief Historical Perspective
Use of mechanical energy to grind down various materials dates back to
the beginning of human history. research activity in the field of mecha-
nochemical processing (McP) has a long history with the first publication
dating back to 1892 when an American chemist M. carrey lea has
shown that the halides of gold, silver, platinum and mercury decomposed
to halogen and the metal during fine grinding in a mor tar [1]. this
study clearly established that chemical changes could be brought about
not only by heating but also by mechanical action. No less reason to
consider M. Faraday, who studied the acceleration of dehydration of the
crystalline hydrates during mechanical action, the founder of mecha no-
chemistry. But the use of mechanically acti vated processes, however,
dates back to the early history of mankind, when fires were initiated by
rubbing flints against one another. W. ost wald coined the term ‘mechano-
chemistry’ in 1891 in the ‘textbook of General chemistry’, which, in
particular, considered various types of stimulation of chemical processes.
While the scien tific basis un derlying McP was investigated from the
very beginning, appli cations of McP products were slow to come about,
mostly because of limi tations on the productivity of McP reactors, pu-
rity of the products and the economics of the process.
It is now accepted that the McP technique embraces three diffe rent
processes, mechanical alloying, mechanical milling, and reaction mil ling.
All these three processes involve cold welding, fracturing and rewelding
of powder particles during repeated collisions with grinding balls in a
high energy milling device. however, depending on the actual process,
other features may be present.
So far, an interest in research in the field of mechanochemistry is
high since the mechanochemical processing is a rather simple and effec-
tive technique for obtaining a wide class of compounds and novel nano-
composite materials (NcM). For the moment, scientists con cent rate
their efforts not only on the technological features and funda mental
principles of McP process, but also on the establishment of application
fields of materials obtained as a result of this synthesis method.
reaction at milling. the efficiency of using carbon nanotubes in the fabrication of
nanocomposite materials with improved functional characteristics is shown. As
revealed, the reaction milling is effective for the synthesis of multicomponent
carbides (substitutional solid solutions).
Keywords: mechanochemical processing, carbon nanotube, carbide, solid solution,
x-ray diffraction, electron microscopy.
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 7
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
1.2. Features of the Mechanical Processes at MCP
the processes taking place at McP are initiated by the mechanical action
of the reactor equipment (ball mill) on the test substance. they belong
to the non-equilibrium phase transformations and occur at room
temperatures, in which the redistribution of atomic compo nents by the
normal diffusion mechanism is absent. however, since the mechanical
action usually causes intense cold plastic defor mation of a substance,
then an additional point or linear defects appear in this substance. the
presence of these defects creates conditions for the transport of atoms
on a distance much larger than interatomic due to the process of strain-
induced mass transfer.
It is known that the temperature is one of the parameters deter-
mining the mobility of atoms in metals and alloys under normal con-
ditions. Increase of the temperature leads to an increase in the diffusion
mobility of atoms. Such factors as radiation, which leads to an increase
of the number of point defects [4, 5], phase (martensitic) transformations
[6–9], and plastic deformation of substance [10–13] resulting in a
growth of the diffusion mobility of atoms by several orders [14] are
necessary to note among other ways to increase the rate of diffusion
processes.
the transfer of large amount of energy accompanies an intense de-
for mation action on a substance in a ball mill, resulting in the formation
of special locally heterogeneous states that are caused by saturation of
the substance by defects and high tensions on sub-micron and nanoscaled
structure elements. Formation of such local stressed states in a substance
results in the two processes, namely: the diffusionless collective dis-
placement of atoms, and the process of anomalous low-temperature dif-
fusion of disordered displacement of atoms at a distance much larger
than interatomic [11–13, 15–17].
Since mechanical treatment of a substance in a ball mill at McP is
purely pulsed process, the mechanical processes described above, as well
as the chemical processes described below, do not occur at the whole
time; but only at the moment of impact and during the period of re-
laxation of the stress field, that takes place on different channels
depending on the synthesis conditions [18].
1.3. Features of the Chemical Processes at MCP
the mechanical processing in ball mills is the most common and quite
simple procedure in mechanochemistry. that is why both the mechanics
and the physics of the processes occurring at grinding are the subject of
various studies, most of which are devoted to optimizing the milling
stage in order to obtain the maximum reaction surface of a substance
with the minimal energy consumption.
8 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
two principles already mentioned are the basis of the research of
processes occurring at mechanochemical processing, namely: the im-
pulse nature of process (the emergence of stress field and its relaxation
[19–22]) and the local character of the mechanical action on the sub-
stance (at processing, the stress field does not occur in the entire vo lu-
me of a substance, but only in the zone of particles contact [23–25]).
the aim of such studies is to determine the boundaries of a region where
the stress field is formed at mechanical action, existence time and a
mode of this region (for example, jump in pres sure [26, 27]), as well as
the determination of relaxation channels of the stress field. In general,
relaxation of the stress field can lead to an increase in the local tem pe-
rature, to the formation of new meta stable phases, to poly morphic
transformations in the substance, to the formation of struc tural defects,
to the emergence of a new active surface, etc. [28].
If the substance that is mechanochemically treated is one-compo-
nent or single-phased the relaxation of a stress field can be accom panied
by the heat release, the formation of metastable and poly morphic phases,
the formation of a new developed particle sur face, the appearance of
defects in the substance, and also by its amorphization.
Part of each of these relaxation channels depends on the conditions
and magnitude of the mechanical load (load rate and an amount of ener-
gy applied), the physical properties of a substance treated, pro cessing
temperature, etc. Sometimes the relaxation channel may change in the
process. For example, an increase in the growth rate of the main cracks
of oxysalts crystals at McP is accompanied by a change in the mechanism
at the tip of crack, i.e., thermal decomposition is replaced by mecha no-
chemical one [29, 30]. change in the size of the particles that are me-
chanically processed is accompanied by the tran sition from their crushing
to the process of plastic deformation [31, 32]. the grinding of a substance
at McP leads to obtaining a maximum reaction surface powder at mini-
mum energy costs that reduces the activation energy of the subsequent
chemical trans for mation along with the accumulation of energy in the
form of defects.
Generally, there are two types of mechanical activation process. It
is referred to a process of the first type if a cumulative time of me cha-
nical action, formation of stress field and its relaxation is longer than
a time of chemical reaction (mechanochemical process). on the contrary,
in the course of a process of the second type, the time of mechanical ac-
tion and the formation of stress field is shorter than the time of chemi-
cal reaction, or in general, these processes are sepa ra ted in time (the
process of mechanical activation or reaction milling).
At reaction milling due to plastic deformation of a substance
processed the accumulated energy is usually consumed to the formation
of defects in activated crystals: dislocations, atomic and ionic vacancies.
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 9
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
In addition, a significant deformation of the crystal lattice of initial
material can lead to the destruction of interatomic bonds, which results
in the formation of free radicals in covalent crystals, and to amorphization
of molecular crystals.
If a multicomponent powder blend is mechanically treated then the
supersaturated solid solutions and stable or metastable inorganic com-
pounds [33–36] may form as a result of solid state mechanochemical
reactions between initial components (mechanical alloying). In addition,
the main factors that determine the possibility of solid solutions forma-
tion by mechanical alloying are the similarity of the atom sizes of com-
ponents and correspondence of their crystalline structures.
It was shown that McP proceeds by a diffusion mechanism, but un-
like the normal diffusion process, it carries with abnormally high values
of diffusion coefficient of the atomic components. this type of diffusion
has been termed ‘deformation atomic entanglement’ or ‘ballistic diffu-
sion’. Its mechanism is significantly different from the mechanism of
normal diffusion, which is determined by the gradients of component
concentrations. there is still no single thought about the mechanism of
Fig. 1. Schematic representation of the stress fields formation factors arising at reac-
tion milling the main channels of their relaxation
10 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
deformation atomic entanglement. It is assumed either that the diffusion
at mechanochemical synthesis is carried out on the interstitial positions
in the lattice [37] or that the channels of this diffusion are dislocations,
the amount of which in the case of McP is constantly increasing [38].
effect of the above factors on formation of the stress field, as well as
the main ways of their relaxation, are schematically depicted in Fig. 1.
In this case, the left side of scheme visualizes the processes of grinding,
and the right part characterizes the reaction milling.
thus, the look back review of processes that are implemented in
mecha nochemical processing and the mechanisms governing these pro-
cesses, revealed the promising use of mechanochemical method for low
temperature synthesis of high-temperature carbide phases. In this paper,
we present experimental data on the McP synthesis of two-component
transition metal carbides in a high-energy plane tary ball mill using the
carbon nanotubes as a carbon component of the charge.
2. Materials and Methods
2.1. Source Materials
charge of the required composition of the initial metals and multiwall
carbon nanotubes (cNt) was used in the mechanochemical synthesis of
carbides and composite materials studied. d-transition metals employed
in a synthesis are summarized in table 1 (purity of metals is no less
than 99.95% by weight). Fig. 2 shows as an example the morphology of
the iron and copper powders used. Micrographs ob tained by scanning
electron microscopy (SeM) demonstrate the uni form size distribution
of metal particles and an absence of coalescent agglomerates. SeM
images of other initial metals are similar to those presented at Fig. 2.
Besides, the calibrated metal filings were used as initial metal materials
for reaction milling of the Me–cNt carbides (Me = V, y, Zr, hf, ta),
table 1.
the multiwalled carbon nanotubes used in this study as a carbon
compo nent of the charge were synthesized by catalytic chemical vapor
deposition method (cVD) at tM Spetzmash ltd (Kyiv, Ukraine). Al2o3,
Moo3 and Fe2o3 oxides were used as the catalysts for the cNt production.
Propylene obtai ned by dehydration of isopropyl alcohol was a source of
the carbon. Parameters of cNts are as follows: the average diameter is
10–20 nm, the specific surface area (determined by argon desorption
method) is 200–400 m2/g and their poured bulk density varies from 20
to 40 g/dm3. teM image of carbon nanotubes is shown in Fig. 3. It is
possible to observe the presence of cNt agglomerates, which essentially
complicates the process of manu fac turing composites with a uniform
cNt distribution by traditional methods of powder metallurgy.
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 11
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
Fig. 2. SeM micrographs of the source powders: a — iron (≈200 µm of size) and
b — copper (≈60 µm of size)
2.2. Mechanochemical Processing
elemental metal powders (or metal filings) and multiwall cNts were
mixed to give the desired average composition and sealed in a vial
(height of 70 mm, diameter of 50 mm) under an argon atmosphere. the
high energy planetary ball mill used for McP is a custom made model
developed at the Metal Physics and ceramics laboratory of the taras
Shevchenko National University of Kyiv. hardened stain less steel balls
(11 units of 15 mm diameter) with a ball-to-powder weight ratio of 40 : 1
were used. the vial temperature was held below 375 K during the
experiments by air cooling. the milling process was cyclic with 15 min
of treatment and 30 min of cooling time. the rotation speed was equal
to 1480 rpm; the acceleration was about 50 g; the pressure for a substance
particle reached 5 GPa.
2.3. X-Ray Diffraction Methods
the full complex of the x-ray diffraction methods (XrD) has been used
to study the kinetics of phase transformations of the initial charge at
mechanochemical processing in a ball mill as well as struc tural changes
in phase components (lattice periods, presence of va can cies, etc.), and
the parameters of real structure (crystallite size, deformation of the
crystalline lattice) of the synthesized carbides.
the XrD data were collected with DroN-3 or DroN-4 automatic
diffractometers (cuKα or coKα radiation, respectively) for the proof
samples selected after a certain milling time. the diffraction patterns
have obtained in a discrete mode under the following scanning parameters:
observation range 2θ = 20–130°, step scan of 0.05°, and counting time
per step at 3 s.
the original software package developed by us for the automated
DroN equipment has been used for analysis and interpretation of the
12 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
x-ray diffraction data obtained. this
package contains a full range of standard
ritveld analysis and is intended for
solving dif ferent XrD tasks, namely,
determination of both peak positions
and integral inten sities of the Bragg re-
f lections by means of full profile ana-
lysis; carrying out qualitative and quan-
titative phase analy sis using PDF data
for phase identifi ca tion and the least square method for lattice constants
refinement; testing of the structure models and refining crystal structure
parameters (in cluding coordinates, atomic position filling, texture,
etc.). the mathe matical algorithms realized for these calculations are
Fig. 3. teM micrographs of the multiwalled
carbon nanotubes
Table 1. Crystallographic data of phases formed in Me–CNT charge
after 60 min of processing in a ball mill
Ме
Initial
material
Structure
type
lattice parameter, nm
RB
cNt
content
calculated,
at.%
Product synthesized Initial metal
а с а с
3d
ti Powder,
90–125 µm
Mg 0.2954 (2) 0.4685 (3) 0.2951 (2) 0.4686 (2) 0.112 ?
V Filings,
<200 µm
α-Fe 0.3027 (3) 0.3030 (3) 0.029 16 (2)
Fe Powder,
<200 µm
α-Fe 0.2866 (2) 0.28665 (4) 0.0029 9 (2)
co Powder,
<80 µm
Mg 0.2508 (2) 0.4076 (8) 0.2507 (2) 0.40695 (4) 0.110 ?
Ni Powder,
<80 µm
cu 0.35225 (2) 0.3524 (1) 0.012 5 (1)
cu Powder,
<60 µm
cu 0.3638 (2) 0.36149 (3) 0.024 2 (1)
4d
y In pieces Mg 0.3641 (2) 0.5747 (5) 0.36474 (3) 0.57306 (4) 0.092 ?
Zr Filings,
<200 µm
Mg 0.3239 (2) 0.5148 (2) 0.3232 (1) 0.5147 (1) 0.098 ?
Nb Powder,
<40 µm
α-Fe 0.33043 (2) 0.3300 (1) 0.013 10 (2)
Mo Powder,
<40 µm
α-Fe 0.3147 (3) 0.3147 (1) 5 (1)
5d
hf In pieces Mg 0.3190 (4) 0.5044 (5) 0.3196 (2) 0.5051 (1) 0.086 ?
ta Filings,
<200 µm
α-Fe 0.3310 (2) 0.33013 (3) 0.025 4 (1)
W Powder,
<40 µm
α-Fe 0.3168 (3) 0.31652 (3) 3 (1)
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 13
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
similar to those used in WincSD software [39]. More information about
this package is available in ref. [40].
typical kinetic curve of the tic phase transformations at MPc is
shown in Fig. 4. It should be noted that test samples selected after 60 mi-
nutes the processing contain ti and tic phases along with a certain
amount of iron (Fig. 4) that appears in the sample due to the wear debris.
Similar curves that visualize the course of the synthesis process
were obtained for each carbide phase synthesized.
crystallites size and deformation of the crystal lattice of the pha ses
synthesized were determined through the Williamson–hall method [41].
the average values of the grain size D and deformation of the crystal
lattice ε of the carbides synthesized were estimated via the peak broa-
dening. the Williamson–hall graphs are plotted as de pen dencies of scaled
broadening of Bragg’s reflections, b* (2θ) = β (2θ) (cos θ)/λ, on scat tering
vector S = (2sin θ)/λ for each test sample have been studied, where θ is
the Bragg’s reflection, λ is the wavelength, β (2θ) = (FWHM2
exp – FWHM2
R)
1/2
is an intrinsic broadening (where FWHMexp and FWHMR are experimental
and instrumental broadening, res pec tively).
typical Williamson–hall graphs plotted for the phases present in
the Fe-tic composite (see section 4.1.1) are presented in Fig. 5. the
average grain size D could be found by extrapolating the b* (2θ) depen-
dencies onto S = 0 axes as D = 1/b* (2θ) at θ = 0. the average micro-
deformation of crystal lattice ε could be found from slope of the b* (2θ)
straight line versus S as ε = b* (2θ)/2S.
taking into account the fact that ε (S) dependence obtained for the
tic phase is a horizontal line (Fig. 5) one can conclude that only a fine
structure of this carbide (and not a deformation of its crystalline lattice)
makes an influence on the diffraction peaks broadening of this phase.
Fig. 5. the Williamson–hall plots for phases of the Fe–tic samples: (1) α-Fe (120
min of processing), (2) tic (120 min of processing), (3) tic (compacted sample), (4)
α-Fe (compacted sample)
Fig. 4. Phase composition (wt.%) of the ti–cNt charge reaction milling vs. the pro-
cessing time
14 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
2.4. Electron Microscopy
the method of scanning electron microscopy (SeM) was used both to
analyze the microstructure of the synthesis products and to determine
the elemental composition of the phase components. SeM exanimation
of the samples was carried out using Jeol JAMP-9500F field emission
auger microprobe operated at 10 kV or a scanning electron microscope
ZeISS eVo 50XVP operated at 15 kV, which offers the flexibility of
optional analysis functions such as an energy-dispersive x-ray spec-
troscopy (eDS).
the detailed analysis of samples microstructure was carried out by
transmission electron microscopy (teM). teM images of the cNts and
composites after milling were obtained with a transmission electron
microscope SelMI PeM-125K operated at 100 kV.
3. characteristic of Interaction
in the Metal–cNT Binary systems
Materials with carbon content are widely used in industry mainly as the
iron based alloyed solid solutions (steels). carbides of transition metals
keen demand as wear-resistant coatings and fillers of metal matrices
when creating solid wear-resistant materials (solid alloys, etc.). tra-
ditionally, the graphite or carbon black are used as a carbon com ponent
for the manufacture of these materials. however, the carbon nanotubes
(cNts), which are characterized by unique combination of the mechanical
characteristics due to their size, geometry and crystal structure [42–44]
look as a promising component of novel functional materials with im-
proved mechanical, thermal, and magnetic properties. carbon nano tubes
form agglomerates in the initial state, which usually do not undergo
destruction and grinding that essentially complicates their use in the
development of materials. however, mechanochemical processing of the
cNt-containing charge in a high-energy ball mill allows us to produce a
variety of commercially useful and scientifically interesting materials
since cNts are amorphized under processing and homogeneously dis tri-
buted in volume. Due to the substantial grinding of the charge and the
destruction of carbon nanotubes, such processing ensures high reactivity
of the components. the interaction of the metal with cNts that occurs
at the same time leads to the formation of supersaturated interstitial
solid solutions and/or carbide phases, usually with a modified structure
and improved properties.
the authors of this review have achieved a certain success in syn-
thesis of the carbide phases by mechanochemical processing of the Me–
cNt charge in a high-energy ball mill [45–47] and these results are
discussed herein.
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 15
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
3.1. Solubility of the Carbon Nanotubes in d-Metals
Studying kinetics of the transition metal carbidization, a special atten-
tion was paid to identify the changes in the initial metals occurring in
their interaction with carbon nanotubes during mechano hemical process.
For this purpose the test samples selected after 60–90 min of processing
of the initial charge in a ball mill have thoroughly tested under XrD
method.
According to the XrD data, at the initial stages of mechanochemical
processing the diffraction patterns from the milling products are similar
to those form the initial metals. Indeed, Figs. 6, 7 present diffraction
patterns for three test samples, representing three different type struc-
tures for the metals studied, namely, the Fe–cNt (α-Fe-type structure),
cu–cNt (cu) and Zr–cNt (Mg). It should be noted that according to the
XrD results the calculated lattice parameters for main phases existing
in the Me–cNt blends practically correspond to those for initial metals
(table 1) after 60 min of processing.
the crystal structure calculations for the milling products selected
after 60 min of processing have revealed that the best agreement between
experimental and calculated intensities of reflections (RB factors are
less than 0.03, table 1) can be achieved for a trial model in which atoms
of carbon are implanting in octahedral pores of cu and α-Fe-type struc-
tures with simultaneous appearance of vacancies in the metal atom
positions. More specifically, the following model is proposed for f.c.c.-
metals: the Fm3m space group, Me atoms are placed in 4 (a) (0 0 0); c
atoms are placed in 4 (b) (0.5 0.5 0.5) (the occupation of 4 (a) position is
slightly less than 1, and 4 (b) position is only partially filled with carbon
atoms). A model proposed for b.c.c.-metals is as follows: the Im3m space
group, Me atoms are placed in 2 (a) (0 0 0); c atoms are placed in 6 (b)
Fig. 7. XrD pattern of the Zr–cNt charge after 60 min of the processing, cuKα
radiation
Fig. 6. XrD patterns of the Fe–cNt and cu–cNt powders treated for 60 min in a
ball mill, coKα radiation
16 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
(0.5 0 0). Similarly to f.c.c.-
structures, the occupation of
2 (a) position is slightly less
than 1, and position 6 (b) is
only partially filled with car-
bon atoms. thus, the cubic
phases studied are Mecx in-
ter stitial solid solutions with
implanting the carbon atoms
into lattices of the initial
metal. Mappings of two typical
structures of cubic cucx and
Fecx interstitial solid solutions
of are shown in Fig. 8. the
calculated amounts of carbon,
which is implanting in the octahedral pores of structures, are listed in
table 1. In the case of phases with hexagonal crystal structures,
calculations made in the Mg-type structure did not lead to a correct
result even taking into account the existing of texture (RB factor was
higher than 0.09, table 1). therefore, it was suggested that the internal
rhombic deformation of the Mg type structure is inherent to these
phases. thus, further cal culations were carried out within the framework
of an orthorhombic lattice with aromb = ahex, romb hex 3b a= , cromb = chex,
Cmcm space group Me atoms are placed in 4 (c) (0 z 0.25); c atoms are
placed in 4 (a) (0 0 0). Similarly to the above cubic structures, occupation
of 4 (c) position is slightly less than 1 while 4 (a) position is only partially
filled with carbon atoms. calculated values of the crystallographic
Fig. 8. Structure mappings of in-
terstitial solid solutions on the XY
axis (see also this figure at the
web-site of the journal)
Table 2. Crystallographic parameters of interstitial
solid solutions on the base of hexagonal metals
Me
lattice parameter, nm
z RB
cNt
content
calcu-
lated,
at.%
Mg type structure Zrcx type structure
a c a b c
ti 0.2954 (2) 0.4685 (3) 0.2945 (3) 0.5111 (5) 0.4687 (3) 0.303 (2) 0.051 8 (2)
co 0.2508 (2) 0.4076 (8) 0.2498 (2) 0.4342 (3) 0.4063 (4) 0.333 (3) 0.056 4 (1)
y 0.3641 (2) 0.5747 (5) 0.3633 (2) 0.6320 (5) 0.5741 (5) 0.318 (5) 0.058 12 (1)
Zr 0.3239 (2) 0.5148 (2) 0.3242 (4) 0.5590 (6) 0.5139 (5) 0.318 (4) 0.055 7 (2)
hf 0.3190 (4) 0.5044 (5) 0.3209 (5) 0.5538 (6) 0.5051 (5) 0.314 (5) 0.051 1
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 17
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
characteristics of the distorted hexagonal phases taking into account
the texture are contained in table 2. It should be noted that this
structural model, describing the above interstitial solid solutions, could
be considered as a new Zrcx type structure.
therefore, performed calculations have shown that the orthorhombic
model better describes the crystal structures of the interstitial solid solu-
tions formed on the base of initial metals with Mg-type structure. the
projection of this structure on the XY plane is shown at Fig. 8. It is seen
that the carbon atoms have an octahedral environment of metal atoms.
thus, XrD study showed that the formation of the interstitial solid
solutions with carbon atoms implanted in the octahedral pores of the
crystal structures of the initial metal takes place at the first stage of
the Me–cNt charge processing in a ball mill. In this case, the metal
sub-lattices of all solid solutions become vacant. Moreover, the lattices
of the metals with Mg-type structure become essentially internally
deformed.
It is natural to assume that the main factor regulating the formation
of these interstitial solid solutions is the diffusion of the carbon atoms
into a metal lattice. Indeed, the analysis provided has shown that the
amount of cNts, which are implanting in the crystal lattice after 60
min of milling (table 3), correlates well with the values of the activation
energy of carbon diffusion in the corresponding metals independently
on the crystal structure of the metal matrix, (Fig. 9).
Table 3. Solubility of the carbon in solid solutions obtained by MCP
and solubility of the carbon in metals at high temperatures
Ме
carbon Solubility, at.%
Activation energy
of diffusion,
kJ ⋅ mole−1
calculated carbon content
in materials studied
Maximum carbon
solubility in the equilib-
rium diagram
3d
ti 8 (2) 3.1 139
V 16 (2) 2.6 113
Fe 9 (2) 0.02 134
co 4 (1) 4.3 155
Ni 5 (1) 2.7 160
cu 2 (1) 0.03 188
4d
y 12 (1) 8.8 ?
Zr 7 (2) 2.0 147
Nb 10 (2) 0.26 134
Mo 5 (1) 1.1 153
5d
hf 1 2.0 209
ta 4 (1) 0.1–7.5 161
W 3 (1) 0.7 169
18 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
According to the data obtained it is clear that if in the state of
thermodynamic equilibrium, the solubility of carbon in the initial
d-metals at high temperatures (above 1500 °с) usually does not exceed
3–4 at.%, than the processing of the Me–cNt charge in a ball mill leads
to the formation of the carbon-rich metastable solid solutions. the
solubility of carbon in these supersaturated solid solutions is usually
substantially higher than its solubility in a stable state (Fig. 9). During
formation of supersaturated solid solutions, the amorphous carbon is
implanting in the octahedral pores of structures of known α-Fe and cu
types, as well as in the octahedral pores of a rhombic distorted Mg type
structures.
3.1.1. Carbides of ІІІb, ІVb and Vb Groups
Process of formation of ІІІb, ІVb and Vb monocarbides with the Nacl-
type structure as well as carbides existing in the y–cNt system have
studied. type and characteristics of the initial metals used are sum-
marized in table 1.
Synthesis of the TiC, ZrC, HfC, VC, NbC, and TaC Monocarbides. A
set of experiments on the mechanical alloying of the equiatomic Me:cNt
(1:1) mixtures was carried out in order to study the Mec process of
formation (table 1). According to the XrD results Mec phase is the
only constituent of the test samples selected after 210–400 min of
processing in a planetary ball mill (samples do not contain any traces of
initial metals). the refined lattice parameters for monocarbides
synthesized are given in table 4. It should be noted that the Mec lattice
parameters obtained here are significantly lower than those which are
inherent to conventional monocarbides (table 4). Fragments of the
typical XrD patterns obtained for the final milling products are shown
in Fig. 10. Average values of the crystalline size D and lattice deformation
ε were defined (table 5) from the broadening of the Mec peaks on XrD
Fig. 9. the dependence of carbon
content in d-metal based solid solu-
tions on the activation energy of the
c diffusion in a metal. Solid solu-
tions obtained by mechanochemical
synthesis are marked by circles. the
maximum solubility of carbon (from
to the data for corresponding state
diagrams) is marked by triangles
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 19
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
patterns. the calculations have revealed that D values for mechanically
alloyed Mec are equal to 10–30 nm.
It should be noted that according to the phase analysis results the
formation of Mec monocarbides is accompanied by wear debris of steel
vial and balls. therefore, some test samples along with the Mec phase
contain an admixture of the α-Fe phase (marked as ‘x’ at Fig. 1), the
amount of which increases with processing time increasing. In order to
eliminate the samples from the iron contamination, we treated all milled
powders in a 50% solution of hydrochloric acid and 5 times washed in
Table 4. Crystallographic data of MeC monocarbides synthesized
by the reaction milling of Me–CNT charge
carbide
experimental data for Me–cNt charge
reaction milling Annealing at 700 °c after reaction milling
lattice parameters,
a, nm
lattice parameters,
a, nm
RB
Me content,
аt.%
Formula
tic 0.4297(9) 0.4301(1) 0.037 48.3 XrD
48.8 eDS
ti0.97c
Zrc 0.4651(8) 0.4655(1) 0.037 49.6 XrD
49.2 eDS
Zr0.99c
hfc 0.4589(1) 0.4588(1) 0.013 50.0 XrD hfc
Vc 0.4154(5) 0.4156(1) 0.034 47.0 XrD
46.6 eDS
V0.94c
Nbc 0.4447(2) 0.4446(1) 0.022 46.8 XrD
45.3 eDS
Nb0.92c
tac 0.4435(9) 0.4436(2) 0.035 46.7 XrD
47.7 eDS
ta0.94c
carbide
reference data for carbides obtained from Me-graphite charge
reaction milling conventional carbides
lattice parameters,
a, nm
Formula
lattice parame-
ters, a, nm
Me content,
аt.%
Formula
tic 0.4311
0.4323
tic
tic0.88
0.4327 51.3 tic0.95
Zrc 0.4651
0.4695
Zrcx Zrc0.83 0.4693 51.3 Zrc0.95
hfc 0.4626 hfcx 0.4642 50.0 hfc
Vc 0.4133
0.4416
Vcx
Vc
0.4167 53.1 Vc0.88
Nbc 0.4422
0.4468
Nbcx Nbc0.91 0.4470 51.3 Nbc0.95
tac – – 0.4456 50.0 tac0.95
20 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
distilled water away from
the Fecl2 precipitate. the
purified powders were dried
at room temperature for
48 hours. According to
eDS data the content of
iron in such powders trea-
ted in hcl does not exceed
0.2 wt.%.
Unfortunately, due to an essential broadening and small number of
diffraction peaks on the diffraction patterns of the final milling products
(Fig. 10), it is impossible to refine the crystal structure of the carbides
formed correctly. consequently, the chemically treated samples have
been annealed for 2 h at 700 °c under argon atmosphere in order to di-
mi nish or even remove additional microdeformations of the crystal
lattice, and thus, to improve the overall appearance of the x-ray patterns
(Fig. 11). XrD study has revealed that the lattice parameters of the
Mec monocarbides existing in the annealed samples are similar to those
in the final milling products (table 4), while diffraction peaks on their
patterns are significantly sharpened (Fig. 11). the D and ε values for
the Mec phases after annealing were determined by the Williamson–
hall plot and approximation method. It should be noted that annealing
is accompanied by a reduction of the lattice strains, while the crystalline
sizes do not change (table 5).
Table 5. Crystallite size and deformation of the crystal lattice
of the MeC carbides formed by the reaction milling of the Me–CNT charge
carbide
Material synthesized After annealing at 700 °c
crystallite size,
D, nm
lattice deformation,
ε, %
crystallite size,
D, nm
lattice deformation,
ε, %
tic 24 (1) 2.54 (5) 22 (2) */11 (1) ** 1.15 (4)
Zrc 33 (3) 1.24 (4) 26 (2)/13 (1) 1.26 (4)
hfc 32 (3) 2.27 (5) 29 (3)/4 (1) 1.59 (5)
Vc 10 (2) 2.34 (5) 12 (2)/7 (1) 1.45 (4)
Nbc 13 (1) 1.39 (4) 8 (1)/8 (1) 0.47 (1)
tac 19 (2) 1.85 (4) 14 (1)/9 (1) 1.73 (2)
* calculated by the Williamson–hall method. ** obtained by the approximation
method.
Fig. 10. XrD patterns of the
milled Me–cNt charge, cuKα
radiation
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 21
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
the XrD results obtained have reviled that the milling time required
for a complete transformation of the raw components into a corresponding
monocarbide correlates with the enthalpy of its formation ΔHf
0 (see
more details later). Moreover, the enthalpy of monocarbide formation
proved to be a suitable parameter to analyse the character of a crystalline
size change in these phases. obviously, the higher the enthalpy of
Fig. 12. the dependences of crystallite size D (a) and concentration of vacancies Cvac
(b) in a metal sublattice on the enthalpy of monocarbide formation
Fig. 11. XrD patterns of the milled Me–cNt charge after treatment by hcl and
annealing at 700 °с, cоKα radiation
22 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
formation, the smaller grain
size of the corresponding Mec
phase (Fig. 12, a).
thus, due to the sharpe-
ning of diffraction peaks it
becomes possible to make a re-
fi nement of the crystal struc-
ture of Mec phases. calcu-
lations required have made in a framework of the Nacl-type structure
(Fm3m space group): Me atoms are placed in 4 (a) (0 0 0) and c atoms
are placed in 4 (b) (0.5 0.5 0.5). In doing so, we have refined both the
occupations of 4 (a) and 4 (b) positions by atoms and the isotropic
temperature factors. the validity of these calculations was confirmed
by the meaning of reliability factor RB that did not exceed 0.04 for each
phase (table 4). results of Mec crystal structure refinements have
revealed the existence of vacancies in 4 (a) position filled with the metal
atoms unlike crystal structure of conventional carbides in which 4 (a)
position is completely filled. In contrast to 4 (a) position, the 4 (b)
position is only partially filled with carbon atoms in both cases. Presence
of vacancies in a metal sublattice results in a significant shift of the
Mec composition to the carbon side, which was also confirmed by eDS
data (table 4).
however, in addition to the proposed model of Nacl-type structure
with 4 (a) and 4 (b) positions partially filled by metal and carbon atoms
the variant of modified Nacl-type structure, in which the partial filling
of 8 (c) (0.25 0.25 0.25) position by additional carbon atoms is realised,
Table 6. Distribution of components by the regular point system
in МеС monocarbides synthesized by MCP
carbide
Distribution of components
RBМе carbon
4 (а) 4 (b) 8 (с)
tic 48.5 48.1 3.4 0.018
Zrc 49.4 48.2 2.4 0.009
hfc 50.0 50.0 0 0.013
Vc 46.8 50.4 2.8 0.014
Nbc 46.1 52.8 1.3 0.015
tac 46.7 52.3 1.2 0.021
Fig. 13. Structure mapping of Vc1+x
carbide on the XY axis (online ver-
sion of the figure is also available)
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 23
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
is possible. the calculations provided have revealed that placing of the
additional carbon atoms in 8 (c) position significantly improves the value
of RB factor, but does not affect the composition of the compound
calculated within the common Nacl-type structure. It is worth noting
that the carbon atoms placed in 8 (c) position have a cubic environment
consisting of metal atoms and other non-equivalent crystallographically
carbon atoms (Fig. 13).
It is possible to calculate the amount of the carbon atoms arranged
in cubic pores assuming that the proposed model of the modified Nacl-
type correctly describes the structure of Mec carbides forming under
McP (table 6).
Using the data of structural calculations, it is possible to estimate
the concentration of vacancies Cvac (%) in Mec monocarbides as Cvac =
= CMe/C′Me, where CMe is the metal content in Mec formed under McP;
C′Me is the metal content in a conventional monocarbide. It is shown that
Cvac values correlates with ΔHf
0 ones for each Mec phase (Fig. 12, b).
Namely, the more complicated the formation of carbide is, the more
vacancies are accumulated in its structure. It is obvious that the presence
of vacancies causes a decrease in a total number of atoms Natom, that has
calculated as for the Mec carbides obtained by McP, as well as for
conventional monocarbides (table 4). It appeared that the calculated
Natom values correlate with the enthalpy of formation for Mec carbides
(Fig. 14, a). Moreover, if this value varies slightly for conventional car-
bides, then this value increases monotonically for reaction milled carbi des
(from Vc to hfc for which Natom corresponds to the stoichiometric one).
In our opinion, the presence of structural vacancies in the metal
sublattice is responsible for a decrease of the lattice parameters of Mec
Fig. 14. the dependences of the calculated value of the total number of atoms in the
structures of Mec carbides synthesized (a) and the partial volume per carbide atom
(b) on the enthalpy of formation (data for conventional Mec are marked as ○, and
data for synthesized products are marked as ●)
24 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al
monocarbides obtained by McP. Using a and Natom values (table 4) one
can calculate the partial volume per one atom in the Mec structure as
V/Natom. the results of calculations have revealed that the values of
partial lattice volume for carbides of the IVb group (tic, Zrc, hfc)
synthesized by McP are very close to those for conventional carbides,
while these values for Vb carbides (Vc, Nbc, tac) are somewhat higher
(Fig. 14, b).
In order to study the morphology of Mec particles formed at McP
of Me–cNt charge, the chemically purified powders were examined by
SeM and eDS methods. As a result, it was shown that all samples do not
contain any raw material traces. According to XrD and eDS data, the
Mec nanocrystals (size of 10–30 nm) form a large powder particles (size
of 10–30 µm, Fig. 15), which are the agglomerates of fine crystals.
Besides, the smaller size of the Mec crystallites (table 5), the smaller
particle is formed (Fig. 15, a). therefore, the largest particles are
inherent to hfc, while the smallest ones are observed for Vc. Moreover,
individual nanocrystalline Mec particles could be found on the surface
of some big particles that is clearly seen in Fig. 15, b, obtained at a
higher magnification.
thus, it was shown that nanoscaled (up to 30 nm) tic, Zrc, hfc,
Vc, Nbc and tac monocarbides with the modified Nacl-type structure
can be successfully synthesized in a short time (80–400 min) from
elemental metals and carbon nanotubes by mechanical alloying (mecha no-
chemical processing) of the charge in a high-energy planetary ball mill.
Fig. 15. SeM micrographs of the milled Me–cNt charge, where magnifications are
×1000 (a) and ×10 000 (b)
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 25
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
Synthesis of the Yttrium Carbides. It is known that the inter action
between yttrium and carbon leads to for mation of a number of other
carbides besides of yc2 carbide. Data on the composition and crystalline
structure of these compounds are quite ambiguous. Particularly, it
concerns y3c (unknown structure), yc0.44 (Nacl-type structure, a =
= 0.5115 nm), y2c (ho2c, a = 0.3617 nm, с = 1.796 nm), y4c5 (Pbcl2, a =
= 0.6574 nm, b = 1.1918 nm, c = 0.3669 nm), y15c19 (Sc15c19, а = 0.794 nm,
с = 1.588 nm), y2c3 (Pu2c3, a = 0.8233 nm), yc2 (cac2, а = 0.3685 nm,
с = 0.6211 nm). exactly this data were used to identify the phases that
are formed in the McP carbides.
According to XrD result, the test sample selected after 60 minutes
of processing in a ball mill does not contain any raw charge materials
(elemental yttrium) (Fig. 16). the two yttrium carbides, namely, a
known cubic yc0.44 carbide with a = 0.5015 nm and a new ycx carbide
whose diffraction pattern was indexing well in a hexagonal lattice with
a = 0.9041 nm, c = 0.6296 nm make a phase composition of this sample.
Further processing of the charge does not lead to a change in the phase
composition of McP products, but lattice parameters of both phases
gradually decrease. thus, for yc0.44 carbide: a = 0.4917 nm after 120 min
of milling and a = 0.4894 nm after 180 min. In turn, for ycx carbide:
a = 0.8978 nm, c = 0.6213 nm after 120 min of treatment and a =
= 0.8903 nm, c = 0.6171 nm after 180 min. Particular emphasis needs
to be placed on that fact that test sample selected after 180 min of
milling contains some amount of α-Fe phase (wear debris). So, one could
jump into conclusion that material synthesized has a high strength since
the abrasion of the vial and balls clearly certifies it.
thus, the experiments provided have shown the efficiency of using
the carbon nanotubes for mechanochemical processing of the yttrium
carbides with enhanced mechanical characteristics.
Fig. 16. Fragment of the
diffractogram of the y–
cNt charge after process-
ing in a ball mill, cuKα ra-
diation. the reflections of
the ycx and yc0.44 carbides
are marked as ‘x’ and ‘O’,
respectively
26 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
3.1.2. Carbides of VIb Group
to study the process of Wc and Mo2c carbides formation a set of
experiments on mechanical alloying of the mixtures containing initial
metal powders and multiwall nanotubes as the carbon component
(Me : cNt ratio is equal to 1 : 1 for Wc and to 2 : 1 for Mo2c) has been
performed. Initial metals used in McP experiments are characterized in
table 1. In order to control the phase transformations occurred during
the charge milling, the phase composition of the test samples selected
after each 1–2 h of processing have studied. According to the XrD
results, all test samples processed up to 4 hours in a ball mill contain no
other phases except an initial metal.
Synthesis of the Mo2C Carbide. According to ref. [62], three poly-
morphous modifications are inherent to Mo2c carbide, namely, high
tem perature (1440–2522 °c) Mo2c with W2c-type structure, high tem-
perature (1200–1440 °c) Mo2c with own type structure (distorted ζ-
Fe2N-type structure) and low temperature (<1200 °c) Mo2c.
the diffraction patterns of the test sample processed for 4 hours
and 5 hours in a ball mill (Fig. 17) are similar to each other and are
Table 7. Crystallographic data for W2C and Mo2C carbides formed
after 4 hours of processing in a ball mill [47]
Atom Site Site occ. x y z
W2c (ζ-Fe2N-type structure)
W 6k 0.93 (1) 0.333 (1) 0 0.280 (3)
c (1) 2d 1.00 (1) 0.333 0.667 0.5
c (2) 1a 1.00 (1) 0 0 0
Space group P31m (no. 162)
lattice parameter, nm a = 0.5168 (3); c = 0.4710 (4)
Independent reflections 27
total isotropic B factor, nm2 B = 1.80 (2) ⋅ 10−2
calculated content, at.% 65.1 (3) W + 34.9 (3) c
reliability factor RB = 0.049
Mo2c (ζ-Fe2N-type structure)
Mo 6k 0.80 (2) 0.333 (1) 0 0.250 (3)
c (1) 2d 1.00 (1) 0.333 0.667 0.5
c (2) 1a 1.00 (1) 0 0 0
Space group P31m (no. 162)
lattice parameters, nm a = 0.5174 (3), c = 0.4744 (4)
Independent reflections 16
total isotropic factor B, nm2 B = 1.92 (6) ⋅ 10−2
calculated content, at.% 61.7 (3) Mo + 38.3 (3) c
reliability factor RI = 0.056
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 27
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
indexing well in a hexa-
gonal lattice with а =
= 0.2991(5) nm, c =
= 0.4744(2) nm. that is
why the trial model for
Mo2c carbide was first
made in the frame of
the simplest W2c-type structure: P3m1 space group, Mo atoms are in
2 (d) (0.333 0.667 0.25) and c atoms are in 1 (a) (0 0 0). While the
refinement of atomic position filling, texture and isotropic temperature
factors providing in this structure model framework led to a good
agreement between experimental and calculated intensities of reflections
at diffraction pattern (the RB reliability factors is about 0.07), the
calculated value of Mo content is equal to 58 (1) at.% Mo, which is too
small for Mo2c phase. So, the structure model of the ζ-Fe2N-type
structure was used as a trial model for Mo2c carbide obtained after 5
hours of processing in a ball mill. It was shown that results of calculation
presented in table 7 characterize the crystal structure correctly [47].
however, the fraction of vacancies in the metal sublattice (≈38 at.% c)
is somewhat higher than those inherent to carbon-rich side (≈35 at.% c)
of solid solution on the base of the high temperature Mo2c carbide. It is
important to note that such a good result was obtained in the case of
calculation in the framework of the orthorhombic Mo2c-type structure
model, as well. thus, a more symmetrical ζ-Fe2N-type structure model
is preferable in our opinion.
As follows from the phase analysis results, the formation of Mo2c
carbide is accompanied with the wear debris (steel vial and balls). Steel
(α-Fe, in fact) interacts with McP products to form a Mo3Fe3c cubic
carbide with a = 1.113 (1) nm.
the average grain sizes D estimated by the classical Williamson–
hall plots are equal to 5–7 nm as well as average relative deformation
of Mo2c lattice ε is equal to 0–0.5%.
Synthesis of the WC Carbide. the W2c carbide was shown to be the
main constituent in the test sample selected after 4 hours of milling
(Fig. 18). According to ref. [62] three polymorphous modi fications are
inherent to this carbide, namely, low temperature 1250–2100 °c W2c
Fig. 17. Fragment of the
diffractogram of the Mo–
cNt charge after process-
ing in a ball mill, cuKα ra-
diation. reflection of the
Mo3Fe3c carbide is indicat-
ed by ‘x’
28 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
with own type structure, high
temperature 2100–2400 °c
W2c with Mo2c-type structure
(distorted ζ-Fe2N-type struc-
ture) and high temperature
2400–2780 °c W2c with ζ-
Fe2N-type structure.
Since the diffraction pat-
terns of phases mentioned abo-
ve are similar to each other (their structures differ in internal deformation
only), each of these type structures was tested as a trial model for
determining the crystal structure of W2c carbide, forming after 4 hours
of processing in a ball mill. First calculation was made in the framework
of the most simple W2c type structure (P3m1 space group, a = 0.2984 (4)
nm, c = 0.4710 (4) nm): W atoms are placed in 2 (d) 0.333 0.667 0.25
positions and c atoms are placed in 1(a) (0 0 0) positions. refining of
atomic position filling, texture and isotropic temperature factors
providing in the framework of this type structure does not lead to a
good agreement between experimental and calculated intensities of
reflections at the diffraction pattern (RB reliability factors does not
exceed the value of 0.09) as well as to anomalous calculated value of
carbide composition (56 (1) at.% W). therefore, further calculations for
the W2c structure were made in the frameworks of the Mo2c-and ζ-Fe2N-
type structures. correctness of these calculations was controlled by the
reliability factors, which do not exceed the value of 0.05 for each model.
that is why a more symmetrical model (ζ-Fe2N-type structure) was
chosen for the W2c carbide structure (table 5). As a result, the calculation
has revealed that vacancies exist in 6 (k) position filled with the tungsten
atoms. Presence of vacancies in a metal sublattice leads to shift of the
W2c carbide’s composition onto ≈35 at.% c, which is inherent to the
carbon-rich side of the solid solution on the base of high temperature
W2c modification.
Further processing of the charge in a ball mill leads to gradual W2c →
→ Wc transformation taking place up to 10 hours of milling (Fig. 19).
As follows from the phase analysis results, the formation of Wc mono-
carbide was accompanied by abrasion of the grinding materials leading
to a formation of the cubic W6Fe6c carbide with a = 1.093(1) nm.
therefore, along with Wc phase the final test sample (10 hours of
milling) contains an admixture of W6Fe6c phase (marked as ‘x’ in
Fig. 18. Structure mapping of Mo2c
and Wc carbides on the XY axis
(see also this figure online)
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 29
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
Fig. 19). crystal struc-
ture calculations revea-
led that Wc carbide
crystallizes into comp-
le tely filled structure
(own type-structure,
space group P6m2, a =
= 0.2874(4) nm, c = = 0.2811 (5) nm): W atoms are located in 1(a) (0 0 0),
while c atoms are occupy positions in 1 (d) (0.333 0.667 0.5) (RB =
= 0.054). the average grain sizes D of the Wc carbide are equal to 4–
7 nm. therefore, the results obtained here have revealed that the
reaction milling of the tungsten and cNts equiatomic mixture results
in a stepwise trans for mation: initial charge → W2c → Wc. At the first
stage (up to 4 hours of processing) the W2c carbide is formed. It should
be noted that a priority of existing the high-temperature W2c carbide
(ζ-Fe2N-type structure) is expected since it is formed congruently at
2775 °c. At the second stage continuing 6 hours, the Wc monocarbide
formation is a result of W2c + cNts → Wc trans formation. Finally, the
Wc monocarbide is the main phase constituent of the product obtained
after 10 hours of the charge processing in a ball mill.
3.1.3. Carbides of VIIIb Group
Process of the Fe3c and co3c carbides formation have studied on test
samples selected stepwise after processing of the initial Me–cNt (3 : 1)
charge in a planetary ball mill by the XrD, SeM and teM methods
Initial metals applied for samples preparation are characterized in table 1.
Synthesis of the Fe3C Carbide According to the XrD data the Fe3c
carbide appears for the first time in the test sample selected after 200 min
of milling the Fe–cNt charge. In this sample Fe3c phase coexists with
a solid solution on the α-Fe base. A significant amount of the Fe3c phase
was detected in a test sample selected after 300 min of milling (Fig. 20).
the indexing of the diffraction pattern obtained (Fig. 20) and
further calculation of crystal structure refinement indicate that the
phase synthesized from the Fe–cNt charge is actually the Fe3c carbide
(table 8). It should be noted that values of the lattice parameters ob-
Fig. 19. Fragments of the
diffractograms of the W–
cNt charge after process-
ing in a ball mill, cuKα
radiation. reflection of the
W6Fe6c carbide is marked
as ‘x’
30 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
tained here for the Fe3c phase are higher than those for conven tional
Fe3c carbide, viz.: a = = 0.5089 nm, b = 0.64343 nm, c = 0.4526 nm.
As a result of crystal structure modeling and refinement (table 8)
it was found that mechano che mical processing of the Fe– cNt charge
leads to the formation of
Fe3c1+x carbide with addi tio-
nal carbon atoms implanting
in the Fe3c crystal lattice. An
increase in the amount of the
carbon atoms leads to a shift
in the com po sition of Fe3c1+x
carbide from 25 to 28 at.% c.
taking into account the exis-
tence of ad ditional atomic po-
sition the crys tal structure
of the Fe3c1+x car bide could
be con sidered as a new type
structure of inor ga nic compo-
unds. the map ping of this car-
bide structure of onto the YZ
plane is pre sented in Fig. 21.
Fig. 21. Structure mapping of the
Fe3c1+x and co3c1+x carbides on the
YZ axis (see also this figure on-
line)
Fig. 20. Fragments of the diffractograms of the Me–cNt charge after processing in
a ball mill, cоKα radiation. reflection of the α-Fe is marked as ‘x’
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 31
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
It is known that the crys tal structure of the Fe3c ce mentite contains
four empty equivalent octahedron with the iron atoms in vertexes (α-Fe6)
and with centres in (0 0 0), (0 1/2 0), (1/2 0 1/2), and (1/2 1/2 1/2).
they are these empty octahedra, which partially filled with carbon atoms
during Fe3c1+x carbide formation (Fig. 21).
Table 8. Crystallographic data for Fe3C and Co3C carbides
formed after 4 hours of processing in a ball mill
Atom Site Site occ. X Y Z
Fe3c (Fec1+x type structure)
Fe (1) 8d 1.00 (1) 0.198 (6) 0.055 (4) 0.322 (6)
Fe (2) 4c 1.00 (1) 0.039 (7) 0.25 0.844 (6)
c (1) 4c 1.00 (1) 0.949 (10) 0.25 0.513 (12)
c (2) 4a 0.17 (4) 0 0 0
Space group Pmna, no. 62
lattice parameters, nm 0.5106 (9), 0.6774 (7), 0.4526 (1)
Independent reflections 62
total isotropic factor B, nm2 B = 2.15 (2) ⋅ 10−2
calculated content, at.% 71.9 (3) Fe + 28.1 (3) c
reliability factor RI = 0.075
Sintered Fe3c (Fec1+x type structure)
Fe (1) 8d 1.00 (1) 0.193 (4) 0.065 (2) 0.348 (5)
Fe (2) 4c 1.00 (1) 0.059 (3) 0.25 0.845 (8)
c (1) 4c 1.00 (1) 0.966 (10) 0.25 0.507 (12)
c (2) 4a 0.32 (8) 0 0 0
Space group Pmna, no. 62
lattice parameters, nm 0.5121 (4), 0.6779 (5), 0.4550 (3)
Independent reflections 62
total isotropic factor B, nm2 B = 3.65 (2) ⋅ 10−2
calculated content, at.% 69.5 (3) Fe + 30.5 (4) c
reliability factor RI = 0.062
co3c (Fec1+x type structure)
co (1) 8d 1.00 (1) 0.209 (2) 0.080 (2) 0.357 (4)
co (2) 4c 1.00 (1) 0.074 (3) 0.25 0.837 (6)
c (1) 4c 1.00 (1) 0.050 (6) 0.25 0.500 (5)
c (2) 4a 1.00 (4) 0 0 0
Space group Pmna, no. 62
lattice parameters, nm 0.4928 (4), 0.6626 (5), 0.4404 (1)
Independent reflections 64
total isotropic factor B, nm2 B = 3.80 (2) ⋅ 10−2
calculated content, at.% 60.0 (4) co + 40.0 (3) c
reliability factor RB = 0.067
32 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
From Fig. 20 representing diffraction patterns of the Fe3c1+x carbide
synthesized, it is evident that all diffraction peaks are broadened
(calculated crystalline size value is equal to 5–8 nm). this fact prevents
to perform crystal structure calculation correctly. that is why the final
powder product of synthesis was sintered at high pressure and high
temperature (hP-ht method, 8 GPa, 850 °c, hol ding time of 40 seconds).
the toroid type-high pressure apparatus was used to create the pressure
(the test powder sample was wrapped into the AlN foil).
the diffraction pattern of the sample consolidated in this way is
shown in Fig. 22. In this bulk sample, the calculated values of lattice
parameters for the Fe3c1+x phase are somewhat higher (a = 0.5121(4)
nm, b = 0.6779 (5) nm, c = 0.4550 (3) nm) than those of the initial
powder product of McP (table 8). the procedure of crystal structure
refinement has confirmed a correctness of the model proposed for the
Fe3c1+x carbide and has shown that the solubility of carbon in this phase
increases to 30.5 at.% under pressure.
Measurements, which were carried out on the consolidated sample
have revealed that average value of the Vickers hardness is equal to
10.4 (3) GPa. this hardness value is similar to that for conventional
Fe3c carbide.
Synthesis of the Co3C Carbide. the x-ray phase analysis of the test
samples selected after a certain processing time of the co–cNt charge
in a ball mill, revealed that the co3c carbide with a diffraction pattern
similar to that for Fe3c (Fig. 20) appears in test samples after 150 min
of milling. the co3c phase coexists with the solid solution on the base
of hexagonal co at earlier milling stages. A significant amount of the
co3c carbide was detected in a test sample selected after 240 min of
processing in a ball mill (Fig. 20).
Fig. 22. Fragment of the
diffractogram of the Fe–
cNt charge consolidated
by hP-ht method, cоKα
radiation
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 33
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
the indexing of the co3c diffraction pattern (Fig. 20) and further
crystal structure refinement within the framework of both Fe3c and
Fe3c1+x type structures has revealed that the variant with additional
carbon atoms is more preferable (table 8). It should be noted that if the
4 (a) position in the Fe3c1+x carbide structure of is only partially occupied
by carbon atoms then the total occupation of the 4 (a) position is
characteristic for co3c1+x carbide. therefore, the composition of this
phase shifts from 25 to 40 at.% c. It is also clear that the implantation
of the additional carbon atoms to the octahedral pores of the co3c
structure should cause a significant increase in the value of the lattice
parameter. however, the comparison of our data for the co3c1+x carbide
(table 8) with the relevant data for conventional co3c carbide (a =
= 0.5033 nm, b = 0.671 nm, c = 0.4483 nm) certifies the likely decreasing
of its lattice parameters at McP. however, the formation of the co3c1+x
carbide is accompanied by a substantial internal deformation of the
initial Fe3c type structure (table 8, Fig. 22). therefore, since the atomic
radius of iron (0.126 nm) is higher than that of cobalt (0.125 nm), the
literature data for co3c is obviously not correct.
Another metastable carbide, namely, co2c (33.3 at.% c) with the
cacl2-type structure, Pnnm space group, a = 0.2910 nm, b = 0.4409 nm,
c = 0.44426 nm was found earlier in the existence range of the co3c1+x
carbide (40 at.% c). however, there were no observed traces of this
carbide in the synthesis product.
thus, the Fe3c1+x and co3c1+x carbides have been synthesized by
mechanochemical processing and their formation is accompanied by the
implantation of additional carbon atoms to the octahedral pores of the
Fe3с type structure.
3.2. Formation Mechanism
of the Transition Metal Carbides at MCP
the eleven carbides of d-transition metals (table 9) were synthesized by
reaction milling of the Me–cNt charge. Note that the metals in a form
of powder or filings as well as in pieces were used as the initial components
of a charge (table 1).
test samples selected after a certain time of processing were studied
mainly by the XrD method. however, SeM and teM methods were also
applied to study the synthesis of Fe3c carbide.
Since the carbon nanotubes are x-ray amorphous, the phase
composition of the synthesis products was additionally controlled by
electron microscopy method. the study has revealed that if the XrD
pattern of a test sample selected after 60 min of processing contains
α-Fe reflections only, then the electron diffraction pattern of this test
sample contains a set of reflections from α-Fe and carbon (Fig. 23, a).
34 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
Table 9. Overall characteristics of carbides synthesized
by MCP of the Me–CNT charge
Initial
metal
carbon content
in a metal
carbide
synthesized
lattice parameters, nm
a b c
Nacl modified type structure
ti 8 (2) tic 0.4301 (1) – –
Zr 7 (2) Zrc 0.4655 (1) – –
hf 1 hfc 0.4588 (1) – –
V 16 (2) Vc 0.4156 (1) – –
Nb 10 (2) Nbc 0.4446 (1) – –
ta 4 (1) tac 0.4436 (2) – –
y 12 (1) ycx 0.4894 (3) – –
ζ-Fe2N type structure
Mo 5 (1) Mo2c 0.5174 (3) – 0.4744 (4)
W 3 (1) W2c 0.5168 (3) – 0.4710 (4)
Wc type structure
W 3 (1) Wc 0.2874 (4) – 0.2811 (5)
Fe3cmod modified type structure
Fe 9 (2) Fe3c 0.5106 (9) 0.6774 (7) 0.4526 (1)
co 4 (1) co3c 0.4928 (4) 0.6626 (5), 0.4404 (1)
Ni 5 (1) – – – –
cu 2 (1) – – – –
Initial
metal
carbon content in a carbide, at.%
enthalpy of formation 0
fH∆ ,
kJ ⋅ mole−1McP conventional technology
Nacl modified type structure
ti 51.5 (5) 48.7 −82.52
Zr 50.6 (4) 48.7 −97
hf 50.0 (3) 50.0 −102.4
V 53.2 (5) 46.9 −57.8
Nb 50.4 (5) 47.7 −66.8
ta 51.3 (4) 48.7 −68.4
y ? ? ?
ζ-Fe2N type structure
Mo 38.3 (5) 35 −5.25
W 35.1 (6) 33 −29
Wc type structure
W 50 (3) 50 −29
Fe3cmod modified type structure
Fe 28 (5) 25 1.6
co 40 (4) 25 0.4
Ni – – –
cu – – –
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 35
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
teM image of this test sample (Fig. 24, a) shows that cNt (gray
threadlike particles) are moving towards to iron particles (black grains)
and partially wrap them. It is known [17] that after 15 min of processing
in a high-energy ball mill, carbon nanotubes are crashed onto the onion-
like particles, and at the further treatment (up to 60 min) cNt are
transformed into amorphous carbon. therefore, it can be assumed that
two processes are realized at the initial stage of the mechanical alloying
of the Fe–cNt charge, namely: amorphization of the nanotubes and the
destruction of the iron particles. together, these processes cause the
penetration of amorphous carbon along the grain boundaries of iron
Fig. 23. Fragment of the diffractogram (cocα radiation) and elec-
tron diffraction patterns of Fe–cNt samples processed in a ball
mill for 60 min (a), 150 min (b) and 240 min (c) [46]
36 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
particles, while the destruction of these boundaries increases the contact
area between iron and carbon, thereby increasing the reactivity at Fe–
cNt interaction.
Further milling of the Fe–cNt charge (longer than 60 min) leads to
the gradual formation of the Fe3c carbide. Although there are no visible
reflections of Fe3c carbide on the diffraction pattern of the test sample
obtained after 150 min of milling, its electron diffraction pattern
contains separate bright spot-like reflections from Fe3c carbide additional
to the diffraction circles inherent to α-Fe (Fig. 23, b). Moreover, there
are no individual carbon nanotube pacticles on the teM image of this
sample (Fig. 24, b). obviously, the amorphous carbon interacts with the
iron atoms forming a certain nucleation center of the Fe3c carbide phase
on the surface of metal particles (the light-colored inclusions on the
surface of the iron particles, Fig. 24, b). Besides, at this stage of the
milling, the iron grains are significantly deformed resulting in the
reducing of the size of their particles.
It is seen (Fig. 24, c) that after 240 min of the Fe–cNt charge
processing of in a ball mill, the process of intensive exfoliation of the
reaction products previously formed on surface of the iron grains begins.
teM image of this sample displays the two types of clusters, namely:
dark spots of iron particles and light stains of Fe3c carbide. In addition,
the electron diffraction pattern of this sample shows a superposition of
the diffracted reflections of iron and Fe3c. Moreover, the presence of
this carbide could be already detected via x-ray diffraction also (Fig. 23).
on this stage, the carbide formation does not completed, but continues
mainly with the participation of carbon atoms, which saturated the
crystal lattice of iron at the initial stages of milling.
obviously, that the above processes of the interaction between the
iron atoms and carbon nanotubes will also take place at the interaction
of cNt with other d-metals under the same technological conditions of
the Me–cNt charge mechanochemical processing.
Fig. 24. тем micrographs of the Fe–cNt samples processed in a ball mill for 60
min (а), 150 min (b) and 240 min (с) [46]
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 37
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
thus, at the first stage of milling (usually up to 60 min of the
charge processing) the total energy of balls collision in a planetary ball
mill is spent mainly on the amorphization of carbon nanotubes and on
crushing the particles of initial metal along their grain boundaries.
together these processes result in an increase of the number of fine
metal particles with a developed surface and enhanced reactivity.
At this stage, amorphous carbon penetrates into the metal lattice,
forming an interstitial solid solution due to the lattice and boundary
diffusion processes (tables 1, 2). During diffusion the carbon atoms
place predominantly in the deformed octahedral pores of the initial met-
al lattices with α-Fe (V, Fe, Nb, Mo, ta, W) and cu (Ni, cu) type struc-
tures, as well as in the octahedral pores of the rhombic distorted lattice
of the initial metals with Mg (ti, co, y, Zr, hf) type structure (Fig. 8).
thus, at the first stage of milling, the main factor, which regulates the
formation of interstitial solid solutions independently on the crystal
structure of the initial metal, is the diffusion of the carbon atoms into
the lattice of corresponding metal. It was shown that the amount of
carbon atoms accumulated by the lattice of metal is well correlated with
the value of the activation energy of the carbon diffusion in the corre-
sponding metal (Fig. 25). As a matter of interest, the metal sublattices
of all solid solutions become vacant to a varying degree. Moreover, the
lattices of metals with Mg type structure become internally deformed.
At the second stage of the charge milling (usually processed from 60
to 250 min), process of the carbon atoms penetration into the metal
matrix is activated. this stage is also characterized by the initiation of
the carbide phase formation on the surface of the initial metal particles.
Generally, the milling time, required for complete transformation of
initial components to the carbide, correlates with the enthalpy of its
Fig. 25. Dependences of
the additional carbon
content in metals and
carbides on the activa-
tion energy of carbon
diffusion, where in-
terstitial solid solutions
are marked as triangles,
carbides are mar ked as
circles
38 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
formation. then, if value of the enthalpy of formation for carbide (tab-
le 9, Fig. 26) is lower (more negative), the process of this carbide
formation is more favorable thermodynamically. talking this into ac-
count Mec monocarbides synthesized could be arranged by the simp li-
city of their formation in a ball mill (from less to longer treatment time)
as hfc → Zrc → tic → tac → Nbc → Vc. Analyzing values of enthalpy
of formations listed in table 9 an approximate time required for a
complete formation of these carbides could be estimated.
A detailed study of the crystal structure of McP phases shows that
the excess carbon content is characteristic both for the carbides and for
interstitial solid solutions. Moreover, the appearance of this excess of
carbon is originated from both the existence of vacancies in metal sub-
lattice of the carbide, and by the implanting of additional carbon atoms
to the empty pores (octahedral or cubic) of the crystal structure of
carbide. the value of the carbon excess in the structure of the McP
carbide (in relation to the carbon content in the structure of the con-
ventional carbide) correlates with the activation energy of carbon diffu-
sion in the corresponding metal (Fig. 24). that is, the two processes are
competing in the formation of carbide phase, namely, diffusion of the
carbon atoms to the lattice of initial metal and the formation of a carbide
phase from a saturated solid solution obtained.
Finally, at the third stage of reaction milling, the formation of
carbide is completed. In this case, the carbide particles are exfoliated
from the surface of metal particles when reaching the corresponding
critical thickness, and subject to milling with a decrease in their
crystalline size.
It is interesting, that the enthalpy value of carbide formation is a
convenient parameter to analyze the character of crystalline size
Fig. 26. Dependence of
the crystallites size of
carbides synthesized at
equal conditions on the
enthalpy of their forma-
tion
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 39
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
variation for McP carbides. Namely, the higher enthalpy of formation,
the faster the carbide is formed.
thus, XrD study, as well as SeM and teM examinations of samples
selected after a certain time of milling of the Me–cNt charge reveals
the most probable mechanism for carbide formation. It is clear that the
mechanical energy of balls collision is transformed into a heating energy,
creating a high-temperature field of local heating. the fact that the
high temperature Mo2c and W2c carbide modifications were formed in
the Mo–cNt and W–cNt systems rather than polymorphic modifications
inherent to these phases at tempe ra tures of synthesis (375 K) argues for
this assumption. local tem perature heating (creation of the temperature
gradients) initiates the process of amorphous carbon diffusion along the
boundaries and inside the grains of the initial metal. enormity of a
diffusion flow results from the supersaturation of solid solutions with
carbon atoms, which place in the octahedral and cubic pores of metal
lattices. Moreover, destruction of this metal lattice due to formation of
structural vacancies also takes responsibility for enhanced diffusion
value. they are an increase in the reactive surface due to crashing of
the metal grains, as well as the presence of local stresses states caused
by the diffusion-induced deformation of lattices of the saturated solid
solutions that create the conditions for formation of just the carbide
phase. therefore, most likely, the fields of mechanical stress are relaxing
in the two main ways, namely, heating and grin ding (Fig. 1). thus, in
this work the carbides of d-transition metals are formed mainly due to
selfsustaining reaction at mechanochemical processing, as previously
was shown for tic in refs. [57, 58].
4. Interaction in the Metal–Metal′–cNT Ternary systems
Metal-matrix composites reinforced with ceramic particles (cermet) have
found their successful application in chemical, aerospace, auto motive,
mining, oil and gas industries due to a unique combination of high wear
resistance, hardness, strength and corrosion resistance. the success in
a synthesis of nanoscale binary transition metal carbides via mechanical
alloying of the Me–cNt charge prompted us to perform a set of expe-
riments on the synthesis of nanocomposite materials (NcM), where the
carbides are formed in the metal matrix directly at milling of the Me–
Me′–cNt charge (Me is the metal matrix, Me′ is the metal component
of a carbide phase). three com positions in the Fe–ti–cNt system and
two compositions in the ti–cu–cNt system were selected as objects of
this study. A number of factors, such as relatively low cost of initial
metals, carbides synthesis rate, simple procedure of consolidation of the
powders milled as well as the expectation of a high strength and plasticity
of NcM obtained gives occasion to this choice.
40 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
4.1. Fe–Ti–CNТ Composites
According to the data for Fe–ti–c isothermal section [64] there
are the three-phase Fe–Fe3c–tic region bordered by the two-phase Fe-
tic and Fe3c–tic regions in a solid state. taking this into account,
three compositions were chosen from these regions: 70 wt.% Fe,
24 wt.% ti and 6 wt.% cNt (56 at.% Fe, 22 at.% ti and 22 at.% cNt)
(charge 1 for Fe–tic region); 80 wt.% Fe, 11 wt.% ti and
9 wt.% cNt (60 at.% Fe, 10 at.% ti and 30 at.% cNt) (charge 2 for
Fe3c–tic region); 72 wt.% Fe, 21 wt.% ti and 7 wt.% cNt)
(56 at.% Fe, 9 at.% ti and 25 at.% cNt) (charge 3 for Fe–Fe3c–tic
region).
According to the XrD results of the test samples, the formation of
tic carbide begins after 30 min of processing of initial charge in a ball
mill. Initial titanium is not detected in the test samples while α-Fe
remains their main constituent after 40 min of milling. Further
processing of charge 1 does not lead to a significant change in its phase
composition, but the milling of both charge 2 and 3 is accom panied by
the appearance and gradual increase of Fe3c phase, which coexists with
the tic carbide. Finally, according to the results of quantitative phase
analysis the test samples selected after 150 min of milling have the
following phase composition:
α-Fe (75) + tic (25) (charge 1),
Fe3c (86) + tic (14) (charge 2),
α-Fe (34) + Fe3c (33) + tic (33) (charge 3)
(Fig. 27), which practically corresponds to their chosen position on the
isothermal section of the Fe–ti–c system.
It should be noted that the primary formation of the tic and not
Fe3c carbide is not surprising since the free energy of formation of the
tic phase is less than that of the Fe3c phase (table 9). that is, the
formation of tic carbide is thermodynamically more favorable than that
of Fe3c carbide. therefore, during McP of the charge the process of tic
formation begins and finishes much earlier than that of Fe3c.
the diffraction peaks on the patterns are significantly broadened
(Fig. 27). therefore, the average values of the grain size D, as well as
the average values of microdeformation of the crystal lattice ε have
determined for all available phases (α-Fe, Fe3c, and tic) using the
Williamson–holl plots. It was shown that the crystalline size is equal to
3–10 nm for all these phases.
the milling process of charge 1 was studied in more details. the
analysis of the experimental data obtained shows that the grain size of
the tic carbide formed during McP is of 3–5 nm. Moreover, for α-Fe
phase the dependences of both the grain size and microdeformation of
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 41
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
lattice on the milling time (Fig. 28) can be approximated by exponential
curves as
0 finexp( )D D t D= − τ + ,
where t is the milling time (min), in the case of grain size curve, D0 and
Dfin parameters can be interpreted as initial and final size (nm) of
crystallites (here, at τ ≈ 20 min, D0 ≈ 1300 nm, and Dfin ≈ 4 nm).
In order to study the mechanical properties of the Fe–ti–cNt
composites, the final powder products, selected after 150 min of milling,
were consolidated by hP-ht sintering. to ensure the desired sintering
parameters (8 GPa, 850 °c, 40 s treatment time), a toroid high pressure
cell was used (the powders were wrapped in AlN foil). According to the
XrD results, the phase compositions of the samples compacted by hP-
ht method are similar to those for powder products (Fig. 28). It is also
shown that microdeformations appear in the crystal lattices of tic
Fig. 27. Fragments of
diffractograms for Fe–
ti–cNt samples conso-
li dated via hP-ht me-
thod (cоKα radiation)
Fig. 28. the processing-time-dependent grain sizes (a) and microdeforma-
tion (b) for the α-Fe phase existing in the Fe–tic composite (charge 1)
42 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
carbide at high pressure sintering, while those inherent to crystal lattice
of α-Fe phase, on the contrary, decreases slightly (stresses in the iron
matrix are partially removed). this fact illustrates Fig. 29, which clearly
shows the sharpening of peaks presented on the diffraction patterns.
the SeM images obtained show homogeneous phase distribution for
the consolidated samples with composition placed on two-phase lines,
namely, on Fe–tic line (charge 1) and on Fe3c–tic line (char ge 2)
(Fig. 30, a, b). It was shown that during hP-ht sintering the nanoparticles
of the final powders (with a grain size of 4–8 nm) processed in a ball
mill are come together in fine-grained (grain size up to 22 nm) solid
materials having a quite high compactness. According to the eDS the
distribution of the elements in these samples is homogeneous and cor-
responds to the composition of char ge 1 and charge 2, respectively.
however, in addition to small inclu sions of the carbide phases the sample
with composition corres pon ding to the Fe–Fe3c–tic three-phase region
(charge 3) contains large separate grains of iron (Fig. 30, c).
It was naturally to assume that the homogeneous and dense distri-
bution of carbide particles in the iron matrix could lead to an increase
Fig. 29. the SeM micrographs of the compacted Fe–ti–cNt samples obtained after
120 min of milling charge 1 (a), charge 2 (b), and charge 3 (c) (×2000)
Fig. 30. XrD patterns of
the samples obtained by
mechanical alloying of
the blend (wt.%): ti : cu
(3 : 1) with 1 vol.% cNt
(powder) followed by the
sintering of this charge
at 980 °c (T denotes ti-
tanium, C — copper, 1 —
ti2cucx)
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 43
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
in the mechanical characteristics of nanocomposites obtained. For all
three NcMs the Vickers’ microhardness has measured on the PMt-3
device at room temperature. Preliminarily all samples have been polished
with a diamond paste. the load of 150 g has been applied to the each
sample for 15 s. Number of indentations per one sample was 50. testing
of the hP-ht sintered composites reveal the high average Vickers’
microhardness values HV: 11.3 (6), 18.3 (4), and 14.6 (3) GPa for the
sample obtained from the charges 1, 2, and 3, respectively.
4.2. Ti–Cu–CNT Nanocomposites
two composition were prepared and processed in the ti–cu–cNt system,
namely, 75 wt.% ti, 25 wt.% cu, (79.9 at.% ti, 20.1 at.% cu) and
67.7 wt.%. ti, 33.3 wt.% cu, (72.7 at.% ti, 27.3 at.% cu) [65]. the 1 vol.%
cNt was added to both of the mixtures. According to the XrD data, 60 min
of the initial charge processing in a ball mill results in the formation of
a nanoscale ti3cu intermetallic (a grain size of about 7 nm). however,
further processing does not lead to a change in the phase composition of
the milling product. Both final products of mechanical synthesis were
compacted by cold pres sing the powders at room temperature and
following sintering at 980 °c. As result of the XrD study, it was found
that the phase compositions of the samples compacted in this way are
absolutely different from those of the powder products of synthesis.
thus, it was shown, that the sample with 25 wt.% cu (20.1 at.%
cu), treated for 60 min in a ball mill and sintered at 980°c, contains a
new phase whose diffraction pattern indexing well in a cubic face-cent-
red lattice with a = 1.1514 (3) nm (this phase is marked as a ti2cucx
carbide in Fig. 21) as well as some addition of the initial titanium [65].
lattice parameter obtained for ti2cucx phase and the characteristic
positions of its diffraction peaks on the diffraction pattern of this
sample give reason to suggest that this phase crystallizes in the ti2Ni-
type structure with implanting carbon atoms. this assumption is quite
natural, as there is well-known ti2cu intermetallic compound, which
crystallizes in this type of structure (ti2Ni one). that is why the
refinement of the crystalline structure of the ti2cucx carbide was carried
out in a model of the ti2Ni-type structure with the tes ting of several
variants of the arrangement of implanting carbon atoms. the most
suitable trial model, which corresponds to the best correlation between
the experimental and calculated values of the intensity of reflections, is
given in table 10. calculation performed in the framework of this model
reveal that the crystal structure of the ti2cuc2 carbide is somewhat
defecting in copper atoms (it contains only 30.8 at.% cu but not 33.3
at.% cu as usually). Besides, it contains a small amount of carbon
atoms implanting therein (only 0.5 at.%). that is, the crystalline
44 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
structure of this carbide is very slightly different from the crystalline
structure of the ti2cu intermetallic that generates it. the projection of
the structure of the ti2cucx carbide on the XY plane is presented in Fig. 31.
however, during the hot temperature sintering, the nanoscale ti3cu
intermetallic existing in the test sample with 27.3 at.% cu, transforms
into a mixture of the above cubic ti2cucx carbide and the ti3cu2cx
phase, whose diffraction pattern are indexing well in a tetra gonal lattice
with a = 1.1985 (2) nm, c = 0.3044 (1) nm (Fig. 32) [65].
Table 10. Crystallographic data for triple carbides formed
at sintering of the milled Ti–Cu–CNT charge [65]
ti2cucx
Atom Site Site occ. X Y Z
ti (1) 48f 1.00 (1) 0.449 (1) 0.125 0.125
ti (2) 16c 1.00 (1) 0 0 0
cu (1) 32d 0.82 (1) 0.209 (1) 0.209 (1) 0.209 (1)
cu (2) 16e 0.16 (1) 0.5 0.5 0.5
c (1) 8a 0.06 (1) 0.125 0.125 0.125
Space group Fd3m, N 227
lattice parameter a, nm 1.1516 (3)
Independent reflections 34
total isotropic factor B, nm2 B = 2.89 (1) ⋅ 10−2
calculated content, at.% 68.6 ti, 30.8 cu, 0.6 c
reliability factor RW = 0.078
ti3cu2cx
атом Position Filling X Y Z
ti (1) 2a 1.00 (1) 0 0 0
ti (2) 8j 0.64 (2) 0.128 (1) 0.182 (1) 0.5
ti (3) 8j 0.76 (2) 0.423 (1) 0.228 (1) 0
ti (4) 4g 1.00 (1) 0.301 (2) 0.199 (2) 0
ti (5) 4h 0.24 (1) 0.250 (1) 0.250 (2) 0.5
cu (1) 4h 1.00 (1) 0.401 (1) 0.099 (1) 0.5
cu (2) 4g 0.64 (1) 0.099 (1) 0.401 (1) 0
cu (3) 8j 0.64 (1) 0.995 (1) 0.206 (1) 0.5
c (1) 4i 0.33 (1) 0 0.5 0.25
Space group P4/mbm, N 127
lattice parameters, a, c, nm 1.1986 (7), 0.3042 (2)
Independent reflections 80
total isotropic factor B, nm2 B = 3.89 (1) ⋅ 10−2
calculated content, at.% 58.3 ti, 37.5 cu, 4.2 c
reliability factor RW = 0.076
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 45
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
Among the phase-analogs
suitable for identification of
the ti3cu2cx phase there is
only ti3cuN nitride with an
unknown crystal structure,
which has close values of the
tetragonal crystal lattice and
similar positions of diffraction
peaks. A verifying of several
trial models calculated in the
P4/mbm space group (preci-
sely this space group has
proposed for the ti3cuN nit-
ride) led to a fully correct model of the crystal structure of the ti3cu2cx
carbide (table 10, Fig. 32). It should be noted that this structure model
could be consi dered as a new, firstly described type structure of the
inorganic compounds. According to the data on the calculated compo-
sition of this compound it is defecting in both titanium and copper
atoms, and it also contains a certain amount of carbon atoms (4.2 at.%),
Fig. 32. XrD patterns
of the samples obtained
by reaction milling of
the blend (wt.%) ti:cu
(2 :1) with 1 vol.% of
cNt (powder) followed
by sintering of this
charge at 980 °c (com-
pact): 1 — ti2cucx, 2 —
ti3cu2cx
Fig. 31. Projection of the carbide
ti3cu2cx structure onto the XY
plane, where blue (black), red
(grey), and green (light) circles de-
note titanium, copper, and carbon
atoms, respectively (for a better
distinguish of different atoms, the
reader is referred to the web ver-
sion of this article)
46 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
which are implanting in the tetrahedral pores of its metal sublattice
(Fig. 31).
thus, the two ti2cucx and ti3cu2cx ternary carbides are formed
as a result of heat treatment (980 °с) of McP products with 20–
27 at.% cu. the crystal structure of the second carbide is determined
for the first time and belongs to a new, previously unknown type
structure.
the microhardness measurements, carried out for compacted samples
with 20.1 and 27.3 at.% cu, revealed that these samples are heterogeneous
in terms of values obtained, which vary within the 6.9–7.1 GPa. that
is, regardless of the phase composition (availability of both ti2cucx and
ti3cu2cx carbides, as well as of other related phases) the average
microhardness of the obtained materials is much higher than that for
pure titanium (0.97 GPa) and for pure copper (0.37). thus, an addition
to the ti–cu mixture with 20–27 at.% cu of a small amount of cNt
makes a significant effect on the mechanical properties of the
nanocomposite materials obtained by compacting the powders treated in
a ball mill.
5. conclusions
A set of experiments on mechanochemical synthesis of d-transition me-
tals binary carbides and composite materials on their basis has performed
using multiwalled carbon nanotubes as a carbon component of the
charge. As a result, nanoscale powders of the eleven carbides were
synthesized, namely, ycx, tic, Zrc, hfc, Vc, Nbc, tac, Mo2c, Wc,
Fe3c, co3c. Besides, most of them were synthesized firstly with the use
of cNt. In addition, nanocomposite materials of the Fe–ti–cNt and
ti–cu–cNt systems were compacted from the powders pro cessed. the
nanocomposite materials obtained possess high hardness value, which is
provided by the formation of a carbide phase in metal matrix at
mechanochemical process directly.
It is shown that at the first stage of mechanochemical processing of
the initial powder blend in a ball mill the crystal lattices of metal
components of the charge are saturated with the carbon amorphized
that leads to formation of interstitial solid solution. At the same time
their metallic sublattice, on the contrary, becomes defecting on me tals.
Both of these factors cause the supersaturation of solid solutions based
on the initial metals. Moreover, the amount of excess carbon (in relation
to its solubility under thermodynamically equilib rium conditions)
correlates well with the activation energy of carbon diffusion into the
lattice of the corresponding metal.
on the second stage the crystal structures of the binary carbides,
which are formed on the basis of distorted lattices of interstitial solid
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 47
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
solution, are usually modified and characterized by additional carbon
atoms, which are placed in the pores of their initial structures (Nacl
and Fe3c type).
the complex of XrD, SeM and teM data have revealed that the
d-transition metal carbides studied in this work at McP are formed
mainly due to a self-sustaining reaction.
efficiency of the carbon nanotubes application for manufacturing of
nanocomposite materials with advanced properties has shown. Mechanical
characteristics of these materials are regulated by the nanoscale metal
matrix and carbide phase, by the high density and by structural features
of their constituents (usually supersaturated with additional carbon).
Generally, the mechanochemical method can be very effective for
the synthesis of multicomponent carbides (substitutional solid solu-
tions), which cannot be obtained in any other ways. Since the duration
of the incubation period (the time before the formation of the carbide
phase begins) correlates with the enthalpy of formation of this carbide,
the criterion for formation of such mutual solid solutions could be the
similarity of their crystal structures and the proximity of the energies
of their formations.
Acknowledgement. the authors appreciate sincerely Prof. V. tkach
and Dr. D. Stra tiichuk (V. N. Bakul Institute for Superhard Material of
the National Academy of Sciences of Ukraine), l. Kapitanchuk (Paton
electric Welding Institute of the National Academy of Science of
Ukraine) and Prof. M. Semen’ko (Department of Physics, taras
Shevchenko National University of Kyiv) for their help in the prepa-
ration of this manuscript and fruitful discussions.
reFereNceS
P. Balaz, 1. Acta Metalurgica Slovaca, No. 4: 23 (2001).
V. Boldyrev, 2. Khimiya v Interesakh Ustoichivogo Razvitiya, 10, Nos. 1–2: 3,
(2002) (in russian).
M. carry-lea, 3. Am. J. Sci., 141: 259 (1891).
e.P. elsukov, I.V. Povstugar, A.l. Ul’yanov, and G.A. Dorofeev, 4. Fiz. Met.
Metalloved., 101, No. 2: 193 (2006) (in russian).
A.c. Damask and G.J. Dienes5. , Point Defects in Metals (Gordon and Breach,
1st edition: 1963).
A.M. Shalaev, 6. Radiatsionno-Stimulirovannaya Diffuziya v Metallakh [radiation-
Induced Diffusion in Metals] (Moscow: Atomizdat: 1972) (in russian).
l.N. larikov and V.M. Kal’chenko, 7. Mekhanizm Vliyaniya Fazovykh Prevra sh cheniy
na Diffuziyu. Diffuziya v Metallah i Splavah [Mechanisms of Influence of the Phase
transformations on Diffusion. Diffusion in Metals] (tula: 1968) (in russian).
S.D. Gercriken and V.M. Fal’chenko, 8. Voprosy Fiziki Metallov i Metallovedeniya,
No. 16: 153 (1962) (in russian).
B.S. Bokshtejn, S.Z. Bokshtein, and A.A. Zhuhovickiy, 9. Termodinamika i Kine-
tika Diffuzii v Tverdykh Telakh [thermodynamics and Kinetics of Diffusion in
Solids] (Moscow: Metallurgiya: 1974) (in russian).
48 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
r.W. Baluffi and A.J. ruoff, 10. J. Appl. Phys., 34, No. 6: 1634 (1963). https://
doi.org/10.1063/1.1702647
A.J. ruoff and r.W. Baluffi, 11. J. Appl. Phys., 34, No. 7: 1848 (1963). https://
doi.org/10.1063/1.172969
V.M. lomer, 12. Vakansii i Tochechnye Defekty [Vacancies and Point Defects]
(Moscow: Metallurgizdat: 1961) (in russian).
yu.P. romashkin, 13. Fiz. Tverd. Tela, 11, No. 12: 1059 (1960) (in russian).
V.V. Neverov, V.N. Burov, and A.I. Korotkov, 14. Fiz. Met. Metalloved., 48, No. 5:
978 (1978) (in russian).
J.S. Benjamin, 15. Sci. Am., 234, No. 5: 40 (1976).
J.S. Benjamin, 16. Mat. Sci. Forum, 88–90: 1 (1992). https://doi.org/10.4028/
www.scientific.net/MSF.88-90.1
P.h. Shingu, 17. Mechanical Alloying, 88–90 (Zurich: trans tech Publ.: 1992).
https://doi.org/10.4028/www.scientific.net/MSF.88-90
V.V. Boldyrev, 18. Eksperimental’nye Metody v Mekhanokhimii Tverdykh Neorga-
nicheskikh Veshchestv [experimental Methods in Mechanochemistry of Solid
Inorganics] (Novosibirsk: Nauka. Siberian branch: 1983) (in russian).
V.V. Boldyrev, 19. Kinetika i Kataliz, 13, 1411 (1972) (in russian).
V. Boldyrev and G. heinicke, 20. Z. Chem. B, 19: 356 (1975).
N.Z. lyahov and V.V. Boldyrev, 21. Izv. SO AN SSSR. Ser. Khim., 5: 8 (1985) (in russian).
N.S. lyakhov, 22. Proc. Second Japan–Soviet Symposium on Mechanochemistry
(eds. G. Jimbo, M. Senna, and y. Kuwohara) (tokyo: Publishing Society Powder
technology: 1988), p. 59.
yu.t. Pavlukhin, ya.ya. Medikov, and V.V, Boldyrev, 23. Izv. SO AN SSSR, Ser.
Khim., 4: 11 (1981) (in russian).
y.t. Pavlukhin, ya.ya. Medikov, and V.V. Boldyrev, 24. J. Solid State Chem., 53,
No. 2: 155 (1984). https://doi.org/10.1016/0022-4596(84)90089-6
yu.t. Pavlukhin, ya.ya. Medikov, and V.V. Boldyrev, 25. Rev. Solid State Sci., 2:
603 (1988).
h. heegn, 26. Proc. First Int. Conf. Mechanochemistry (cambridge: cambridge In-
tersci. Publ.: 1993), p. 11.
r.B. Schwarz and c.c. Koch, 27. Appl. Phys. Lett., 49, No. 3: 146 (1986). https://
doi.org/10.1063/1.97206
D.r. Maurice and t. courtney, 28. Metall. and Mat. Trans. A, 21: 289 (1990).
https://doi.org/10.1007/BF02782409
V.V. Boldyrev, V.r. regel’, o.F, Pozdnyakov, F.h. Urukaev, and B.ya. Byl’skij, 29.
Dokl. AN SSSR, 221: 634 (1975) (in russian).
F.h. Urukaev, V.V. Boldyrev, o.F, Pozdnyakov, and V.r. regel’, 30. Kinetika i Ka-
ta liz, 18, 350 (1977) (in russian).
e.l. Goldberg, S.V. Pavlov, 31. Proc. Second World Congress on Particle Technology
(ed. G. Jimbo) (Kyoto: Japan Society technology: 1990), p. 507.
V.V. Boldyrev, S.V. Pavlov, and e.l. Goldberg, 32. Intern. J. Miner. Proc., 44–45:
181 (1996).
c.c. Koch, 33. Mater. Trans., JIM, 36, No. 2: 85 (1995). https://doi.org/10.2320/
matertrans1989.36.85
c. Suryanarayana, 34. Progress Mater. Sci., 46, Nos. 1–2: 1 (2001). https://doi.
org/10.1016/S0079-6425(99)00010-9
c. Suryanarayana and N. Al-Aqeeli, 35. Progress Mater. Sci., 58, No. 4: 383 (2012).
https://doi.org/10.1016/j.pmatsci.2012.10.001
t.F. Grigorieva, A.P. Barinova, and N.Z. lyakhov, 36. Russ. Chem. Rev., 70: 45
(2001). https://doi.org/10.1070/rc2001v070n01ABeh000598
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 49
Effect of Carbon Nanotubes on Synthesis of Nanopowders and Nanocomposites
P.y. Butyagin, 37. Russian Scientiéc Review. Sect. B: Chemistry Reviews, 2, Pt. 2:
89 (london: harwood Academic Publ.: 1998).
r. Schwarz, 38. Mater. Sci. Forum, 269–272: 665 (1998). https://doi.org/10.4028/
www.scientific.net/MSF.269-272.665
V.K. Pecharsky and P.y. Zavalij, 39. Fundamentals of Powder Diffraction and
Structural Characterization of Materials (New-york: Springer: 2009). https://
doi.org/10.1007/978-0-387-09579-0
M. Dashevskyi, o. Boshko, o. Nakonechna, and N. Belyavina, 40. Metallofiz. No vei-
shie Tekhnol., 39 No. 4: 541 (2017). https://doi.org/10.15407/mfint.39.04.0541
G.K. Williamson and W.h. hall, 41. Acta Met., 1, No. 1: 22 (1953). https://doi.
org/10.1016/0001-6160(53)90006-6
S. Iijima, 42. Nature, 354: 56 (1991). https://doi.org/10.1038/354056a0
X. long, y. Bai, M. Algarni, y. choi, and Q. chen, 43. Mat. Sci. Eng. A, 645: 347
(2015). https://doi.org/10.1016/j.msea.2015.08.012
r.A. Andrievskiy and A.V. ragulya, 44. Nanostrukturnye Materialy [Nanostructured
Materials] (Moscow: Academiya: 2005) (in russian).
o. Boshko, o. Nakonechna, M. Dashevsky, K. Ivanenko, N. Belyavina, and S. revo, 45.
Adv. Powder Technol., 27, No. 4: 1101 (2016). https://doi.org/10.1016/j.apt.
2016.03.019
o. Boshko, o. Nakonechna, N. Belyavina, M. Dashevsky, S. revo, 46. Adv. Powder
Technol., 28, No. 3: 964 (2017). https://doi.org/10.1016/j.apt.2016.12.026
o. Nakonechna, M. Dashevskyi, and N. Belyavina, 47. Metallofiz. Noveishie Tekhnol.,
40, No. 5: 637 (2018). https://doi.org/10.15407/mfint.40.05.0637
P. Matteazzi and G. le ca48. ër, J. Am. Ceram. Soc., 74, No. 6: 1382 (1991). https://
doi.org/10.1111/j.1151-2916.1991.tb04116.x
A. teresiak and h. Kubsch, 49. Nanostruct. Mater., 6, Nos. 5–8: 671 (1995). https://
doi.org/10.1016/0965-9773(95)00147-6
Q. yuan, y. Zheng, and h. yu, 50. Int. J. Refract. Met. Hard Mater., 27, No. 4: 696
(2009). https://doi.org/10.1016/j.ijrmhm.2008.11.003
h. Jia, Z. Zhang, Z. Qi, G. liu, and X. Bian, 51. J. Alloys Compd., 472, Nos. 1–2:
97 (2009). https://doi.org/10.1016/j.jallcom.2008.04.070
B. Ghosh and S.K. Pradhan, 52. Mater. Chem. Phys., 120, Nos. 2–3: 537 (2010).
https://doi.org/10.1016/j.matchemphys.2009.11.048
c.J. lu and Z.Q. li, 53. J. Alloys Compd., 395, Nos. 1–2: 88 (2005). https://doi.org/
10.1016/j.jallcom.2004.11.046
N.J. calos, J.S. Forrester, and G.B. Schaffer, 54. J. Solid State Chem., 158, No. 2:
268 (2001). https://doi.org/10.1006/jssc.2001.9107
l. takacs, 55. J. Solid State Chem., 125, No. 1: 75 (1996). https://doi.org/10.1006/
jssc.1996.0267
B.h. lohse, A. calka, and D. Wexler, 56. J. Alloys Compd., 434–435: 405 (2007).
https://doi.org/10.1016/j.jallcom.2006.08.216
N.Q. Wu, G.X. Wang, J.M. Wu, Z.Z. li, and M.y. yuan, 57. Int. J. Refract. Met.
Hard Mater., 15, Nos. 5–6: 289 (1997). https://doi.org/10.1016/S0263-4368(97)
87504-X
X.K. Zhu, K.y. Zhao, B.c. cheng, Q.S. lin, X.Q. Zhang, t.l. chen, and y.S. Su, 58.
Mater. Sci. Eng. C, 16, Nos. 1–2: 103 (2001). https://doi.org/10.1016/S0928-
4931(01)00283-1
e.P. elsukov and G.A. Dorofeev, 59. Khimiya v Interesakh Ustoichivogo Razvitiya,
10: 59 (2002) (in russian).
e.P. elsukov, G.A, Dorofeev, and V.V. Boldyrev, 60. Khimiya v Interesakh Ustoi-
chivogo Razvitiya, 10: 53 (2002) (in russian).
50 ISSN 1608-1021. Prog. Phys. Met., 2019, Vol. 20, No. 1
O.I. Nakonechna, M.M. Dashevskyi, О.І. Boshko, et al.
e.P. yelsukov, G.A. Dorofeev, V.A. Barinov, t.F. Grigorieva, and V.V. Boldyrev, 61.
Mater. Sci. Forum, 269–272: 151 (1998). https://doi.org/10.4028/www.
scientific.net/MSF.269-272.151
N.P. lyakishev, 62. Fazovye Diagrammy Binarnyh Metallicheskikh Sistem [Phase
Diagrams of Binary Metallic Systems] (Moscow: Mashinostroeniye: 1996) (in
russian).
y.B. li, B.Q. Wei, J. liang, Q.yu, and D.h. Wu, 63. Carbon, 37, No. 3: 493 (1999).
https://doi.org/10.1016/S0008-6223(98)00218-8
V. raghavan, 64. J. Phase Equilib., 24, No. 1: 62 (2003). https://doi.org/10.1007/
s11669-003-0010-8
o.I. Nakonechna, N.N. Belyavina, M.M. Dashevskyi, K.o. Ivanenko, and S.l. re-65.
vo, Phys. Chem. Solid State, 19, No. 2: 179 (2018). https://doi.org/10.15330/
pcss.19.2.179-185
received october 1, 2018;
in final version, January 3, 2019
О.І. Наконечна 1, М.М. Дашевський 1,
О.І. Бошко 2, В.В. Заводянний 3, Н.М. Білявина 1
1 Київський національний університет імені тараса шевченка,
просп. Глушкова, 4, 03022 Київ, україна
2 Інститут металофізики ім. Г.в. Курдюмова нан україни,
бульв. академіка вернадського, 36, 03142 Київ, україна
3 Херсонський державний аграрний університет,
вул. стрітенська, 23, 73006 Херсон, україна
вПЛИв вуГЛецевИХ нанОтрубОК
на меХанОХІмІчнИй сИнтез нанОПОрОшКІв
КарбІДІв d-метаЛІв І нанОКОмПОзИтІв на їХ ОснОвІ
механохімічним методом у високоенергетичному планетарному кульовому мли-
ні з шихти, що містить вуглецеві нанотрубки, синтезовано нанорозмірні моно-
(порошки) та подвійні (компактовані нанокомпозити) карбіди d-перехідних ме-
талів. розглянуто вплив багатошарових вуглецевих нанотрубок на механохімічний
синтез одержаних матеріалів. з’ясовано особливості механізму формування кар-
бідів перехідних металів у процесі механохімічного синтезу. зокрема, пока за но,
що на першому етапі синтезу (до 60 хв оброблення шихти) має місце аморфізація
вуглецевих нанотрубок і подрібнення частинок вихідного металу по межах зерен.
надалі аморфізований вуглець потрапляє всередину ґратниці металу, утворюючи
твердий розчин втілення, внаслідок чого металеві підґратниці набувають де фект-
ности. на другому етапі синтезу (від 60 до 250 хв розмелювання) процес втілення
атомів Карбону в металічну матрицю пришвидшується та починається форму-
вання карбідних фаз на поверхні частинок вихідного металу. третій етап синтезу
завершує формування карбіду. встановлено, що час розмелювання, по трібний
для повного перетворення вихідних компонентів у карбід, корелює з ентальпією
його утворення, а поля механічних напружень релаксують за двома основними
напрямами: нагрівання та подрібнення. вияв лено, що досліджені в даній роботі
карбіди d-перехідних металів під час ме ханохімічного синтезу фор муються в ос-
новному за рахунок самопідтримуваної реакції. Показано ефек тив ність вико-
ристання вуглецевих нанотрубок при ство ренні наноком по зиційних ма теріалів з
поліпшеними функціональними ха рак теристиками. встановлено, що механо хі-
ISSN 1608-1021. Usp. Fiz. Met., 2019, Vol. 20, No. 1 51
Effect of Cyclic Martensitic γ–ε–γ Transformations on Diffusion Characteristics
мічний метод є ефективним для синтезу багатокомпонентних карбідів (твердих
розчинів заміщення), одержати які ін ши ми способами майже неможливо.
Ключові слова: механохімічний синтез, вуглецева нанотрубка, карбід, твердий
розчин, рентгенівська дифракція, електронна мікроскопія.
О.И. Наконечная 1, Н.Н. Дашевский 1,
О.И. Бошко 2, В.В. Заводянный 3, Н.Н. Белявина 1
1 Киевский национальный университет имени тараса шевченко,
просп. Глушкова, 4, 03022 Киев, украина
2 Институт металлофизики им. Г.в. Курдюмова нан украины,
бульв. академика вернадского, 36, 03142 Киев, украина
3 Херсонский государственный аграрный университет,
ул. стретенская, 23, 73006 Херсон, украина
вЛИянИе уГЛерОДныХ нанОтрубОК
на меХанОХИмИчесКИй сИнтез нанОПОрОшКОв
КарбИДОв d-метаЛЛОв И нанОКОмПОзИтОв на ИХ ОснОве
механохимическим методом в высокоэнергетической планетарной шаровой мель-
нице из шихты, содержащей углеродные нанотрубки, синтезированы нанораз-
мерные моно- (порошки) и двойные (компактированные нанокомпозиты) кар-
биды d-переходных металлов. рассмотрено влияние многослойных углеродных
на нотрубок на механохимический синтез полученных материалов. выяснены осо-
бенности механизма формирования карбидов переходных металлов в процессе
механохимического синтеза. в частности, показано, что на первом этапе синтеза
(до 60 мин обработки шихты) имеет место аморфизация углеродных нанотрубок
и измельчение частиц исходного металла по границам зёрен. в дальнейшем
аморфизированный углерод попадает внутрь решётки металла, образуя твёрдый
раствор внедрения, в результате чего металлическая подрешётка становится де-
фектной. на втором этапе синтеза (от 60 до 250 мин размола) процесс внедрения
атомов углерода в металлическую матрицу ускоряется и начинается формирова-
ние карбидных фаз на поверхности частиц исходного металла. третий этап син-
теза завершает формирование карбида. установлено, что время размола, необхо-
димое для полного преобразования исходных компонентов в карбид, корре лирует
с энтальпией его образования, а поля механических напряжений релаксируют
по двум основным каналам: нагрев и измельчение. Обнаружено, что исследован-
ные в данной работе карбиды d-переходных металлов при механохимическом
синтезе формируются в основном за счёт самоподдерживающейся реакции. По-
казана эффективность использования углеродных нанотрубок при создании на-
нокомпозиционных материалов с улучшенными функциональными характери-
стиками. установлено, что механохимический метод является эффективным для
синтеза многокомпонентных карбидов (твёрдых растворов замещения), получить
которые другими способами практически невозможно.
Ключевые слова: механохимический синтез, углеродная нанотрубка, карбид,
твер дый раствор, рентгеновская дифракция, электронная микроскопия.
|