Equiatomic below 5 at.% [8]. Improvements in the mechanical

Equiatomic or near-equiatomic high entropy alloys (HEAs) have attracted much attention due to
their potential beneficial mechanical characteristics, such as high strength,
high fracture toughness, good
ductility and good wear resistance 1.  HEAs containing
five or more elements, with each elemental concentration
between 5 at.%
and 35 at.%, have high configurational entropy (Smix>1.61R where R is the gas constant)
which suppresses the formation of an
intermetallic phase and
may favor the
formation of simple fcc
or bcc structures
based on solid solution phases leading to
the formation of fcc or bcc solid solutions 2,3.

In
order to increase the strength of these alloys without significantly
sacrificing their
ductility, additional strengthening methods may be introduced such as solid solution or
precipitation hardening which
requires thermomechanical
processing or modifications of the chemical
compositions
of the alloys 4-7.
It is well known that, in
addition to the principal elements, HEAs may contain minor
elements with each below 5 at.% 8.  Improvements in the mechanical properties
of metallic alloys is traditionally achieved by doping 9 and doping elements
such as carbon for the formation of high-strength fine carbides may lead to
improvements in the strength of HEAs. In addition, it is reasonable to anticipate that the additional
strengthening method of grain
refinement, achieved through processing by severe plastic deformation (SPD), may
lead to improvements in the
mechanical properties of HEAs. Equal-channel angular pressing (ECAP) 10 and high-pressure
torsion (HPT) 11 are
well-established SPD procedures
for achieving ultrafine and even
nanostructured grains in metals and alloys but HPT is generally advantageous because, by
comparison with ECAP, it produces smaller grain sizes 12,13 and higher
fractions of grain boundaries having high angles of misorientation 14.

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There are now several reports describing the
influence of HPT processing on HEAs 15-24 and a single report documenting the processing of
an HEA by ECAP 25.  Furthermore, it
was shown earlier that processing a
CoCrFeNiMn high-entropy alloy by HPT leads to significant hardening and grain
refinement and, in addition, post-deformation annealing (PDA) of the
nanocrystalline CoCrFeNiMn HEA provides an excellent combination of high
strength and good ductility 19. Experiments also showed that a
nanocrystalline CoCrFeNiMn HEA processed by HPT exhibited excellent ductility
at elevated temperatures including superplastic elongations with a maximum
elongation of >600% at a testing temperature of 973 K 22,26.

One of the potential methods for
increasing the strength of HEAs with an fcc crystal lattice is doping by
carbon for the formation of high-strength fine carbides. Nevertheless, the
influence of these carbides on the structure and properties of HEAs has received
only limited attention.  Accordingly, the
present research was initiated to evaluate the effect of grain refinement due
to HPT on the microstructures of two CrFe2NiMnV0.25 HEAs doped by carbon.

Experimental materials and procedures

CrFe2NiMnV0.25C0.075
and CrFe2NiMnV0.25C0.125 high entropy alloys were prepared using a non-consumable vacuum arc
melting technique in a water-cooled copper crucible. The purities of the
alloying elements were above 99.9 at.%. After several remeltings (5 times) for
homogenization, the ingots were hot forged and then homogenized at 1323 K for 1
h.  Then the samples were cold-rolled to
~30 % thickness reduction followed by annealing at 1373 K for 1 h. To prevent
oxidation, all samples were sealed in vacuum quartz tubes filled with titanium
chips before the heat treatments.  Polished
disks with diameters of 10 mm and thicknesses of ~0.8 mm were prepared from the
annealed alloys and then processed by HPT for 5 turns at room temperature (RT)
under an applied pressure of 6.0 GPa at 1 rpm using quasi-constrained
conditions in which there is a small outflow of material around the peripheries
of the disks during the torsional straining 27.

After processing, each HPT disk was polished to a
mirror-like quality and hardness measurements were taken using a Vickers
microhardness tester with a load of 500 gf and dwell times of 10 s. The average
microhardness values, Hv, were measured along randomly selected diameters on
each disk. These measurements were taken at intervals of ~0.5 mm and at every
point the local value of Hv was obtained from the average of four separate
hardness values. The phase constituents were determined using X-ray diffraction
(XRD) employing Cu K? radiation (wavelength ? = 0.154 nm) at 45 kV and a tube
current of 200 mA with Rigaku SmartLab equipment. The XRD measurements were
performed over 2? angular ranges from 30° to 100° using a scanning step of
0.01° and a scanning speed of 2º min-1. 

Microstructural characterizations were carried out
using optical microscopy (OM) and transmission electron microscopy (TEM). Foils
for TEM were prepared before PDA and after PDA at 823 K for 60 min using a
focused ion beam (FIB) Zeiss Nvision 40 FIB facility at 3 mm from the CrFe2NiMnV0.25C0.125
HEA disk centres in the normal sections of the disks so
that the normals of the images lay in the shear direction. The TEM micrographs
were obtained using a JEOL JEM-3010 microscope operating under an accelerating
voltage of 300 kV.

Experimental results

Initial
microstructure of  HEA

Figure 1 shows the microstructure of the
as-cast CrFe2NiMnV0.25C0.125
HEA. The observations reveal an extended dendritic structure in Fig. 1(a)
including darker dendrites and lighter interdendritic areas as shown in the
higher magnification image in Fig. 1(b). The inter-dendrite spacing was about
several micrometers and the average grain size was measured as ~300 mm. The
microstructure contained approximately 6% of carbides (Cr23C6
or Cr22.23Fe0.77C6) in the form of particles.
The microstructure of the CrFe2NiMnV0.25C0.075 HEA revealed
similar regularities. It was shown earlier for the same alloy that the second phase in the interdendritic eutectic is chromium carbide wherein some of the
chromium atoms are substituted by other elements 28.

Hardness
evaluations before and after HPT processing

Figure
2 shows the results for the Vickers microhardness
measurements after processing through 5 turns with the average values of Hv
plotted along each disk diameter and with the lower dashed lines at Hv ? 170
and 180 corresponding to the initial hardness in the coarse- grained CrFe2NiMnV0.25C0.075
and CrFe2NiMnV0.25C0.125
HEAs, respectively. The results clearly show that
the hardness at the edge of the disk increases significantly after 5 turns by a
factor of ~2.5 to Hv ? 430 and ~435 , with reference to the annealed condition,
for the CrFe2NiMnV0.25C0.075
and CrFe2NiMnV0.25C0.125
HEAs, respectively. These results
demonstrate that after 5 turns it is not possible to produce a
fully-homogeneous hardness distribution and instead the results show there is a
very small area, within a radius of r dislocations on the primary {111} planes with part of the
dislocations splitting into 1/6 Shockley partials
containing stacking faults.  Hard
precipitates in the fcc HEAs, such as particles of the ?-phase,
act as strong barriers for dislocation motion so that most plasticity takes
place in the matrix fcc phase 33. The mechanisms of plasticity in fcc
nanocrystalline materials are not fully understood at the present time but
generally it is considered that grain boundary sliding is especially important
at grain sizes below ~10-15 nm whereas for larger grains the motion of partial
dislocations becomes dominant and this changes to the gliding of full
dislocations when the grain size increases above ~100 nm 34.  Below ~100-200 nm the plasticity is
controlled by the nucleation and annihilation of dislocations at grain
boundaries.  At present no specific
information is available on the plasticity mechanisms in nanocrystalline fcc
HEAs but it is suggested that probably their behaviour is similar to the
mechanisms in conventional fcc nanocrystalline alloys. Accordingly, for
HPT-processed nanocrystalline CrFe2NiMnV0.25C0.075
and CrFe2NiMnV0.25C0.125 HEAs with average
grain sizes of ~25 nm it is reasonable to anticipate that the splitting of
dislocations into partials plays an important role by comparison with their
coarse-grained counterparts.

If the strength of the nanocrystalline materials is
controlled by the nucleation of dislocations at grain boundaries which depends
directly on the grain size, then the other contributions to yield strength
typical for coarse-grained materials, such as hard particles of second phases
and interstitial impurities 35, should be less important in the
nanocrystalline state. Such behaviour can be observed by comparing the
hardening by carbides in the CrFe2NiMnV0.25C0.075
and CrFe2NiMnV0.25C0.125 HEAs in Fig. 2 with
earlier reports for the single phase CoCrFeNiMn HEA 16,19.  In the coarse-grained state the hardening by
carbides leads to higher strength and higher hardness for CrFe2NiMnV0.25C0.125
by comparison with CoCrFeNiMn while in the nanocrystalline state the single
phase CoCrFeNiMn HEA has a higher hardness of ~450 Hv due to the smaller grain
size of ~10 nm 19 compared with an average grain size of ~30 nm and a
hardness of ~435 Hv in the CrFe2NiMnV0.25C0.125
HEA.  It is concluded, therefore, that
hardening by carbides in the CrFe2NiMnV0.25C0.125
HEA is less effective than hardening through the smaller grain size in the
CoCrFeNiMn HEA.

Thermal stability of HEAs after PDA

The microstructures of the two HEAs consist of an fcc
phase and chromium carbide particles before and after HPT processing but
further annealing at 823 K leads to formation of new precipitates.  This is
consistent with the substantial increase in hardness upon annealing and with
the XRD results.  Formation of precipitates in a single phase
HEA after annealing within special temperature ranges is a well-known
phenomenon 36-39. For example, it was shown that the CoCrFeMnNi alloy has a
single-phase fcc structure above 873 K but a mixture of fcc and bcc
phases, or under some conditions a ? phase with a tetragonal crystal
structure, below 873 K 19.  In
addition, CoCrFeNiMnCx (x = 0.1, 0.175, 0.25) HEAs demonstrate a
significant increase in hardness after annealing in the temperature range of
875-1275 K due to the formation of precipitates 31. The present results
confirm the formation of a multi-phase nanostructured HEA after PDA at 823 K
consisting of precipitates distributed within the microstructure.

Close
inspection of XRD results (Fig. 5) reveals the ? phase in the initial annealed and HPT-processed
samples after annealing at 823 K. Nevertheless, these results indicate that the volume fraction of
precipitates in the HPT-processed sample is higher than in the initial annealed
sample. Basically, it is
well known that HEAs have sluggish diffusion which affects diffusion controlled
mechanisms such as precipitating and grain coarsening 40-42. It appears that the
large number of grain boundaries and imposed defects in the nanocrystalline
HPT-processed promote fast diffusion pathways and also as preferential
nucleation sites for the formation of precipitates. Thus, the severe
plastic deformation leads more quickly to the formation of stable precipitates
compared with fully
annealed samples.

Inspection of the hardness results in Table 1 shows
the hardness decreases significantly above 773 K and up to 1273 K and this is due
to dissolution of the precipitates and activation of the grain coarsening. Thus,
the final hardness at 1273 K is almost the same as in the initial annealed
condition due to grain
coarsening. Therefore, the dissolution of the precipitates plays an important
role in the stability of the microstructure and grain coarsening during annealing at 823
K.