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

Equiatomic or near-equiatomic high entropy alloys (HEAs) have attracted much attention due totheir potential beneficial mechanical characteristics, such as high strength,high fracture toughness, goodductility and good wear resistance 1.  HEAs containingfive or more elements, with each elemental concentrationbetween 5 at.%and 35 at.%, have high configurational entropy (Smix>1.61R where R is the gas constant)which suppresses the formation of anintermetallic phase andmay favor theformation of simple fccor bcc structuresbased on solid solution phases leading tothe formation of fcc or bcc solid solutions 2,3.

Inorder to increase the strength of these alloys without significantlysacrificing theirductility, additional strengthening methods may be introduced such as solid solution orprecipitation hardening whichrequires thermomechanicalprocessing or modifications of the chemicalcompositionsof the alloys 4-7.It is well known that, inaddition to the principal elements, HEAs may contain minorelements with each below 5 at.% 8.  Improvements in the mechanical propertiesof metallic alloys is traditionally achieved by doping 9 and doping elementssuch as carbon for the formation of high-strength fine carbides may lead toimprovements in the strength of HEAs. In addition, it is reasonable to anticipate that the additionalstrengthening method of grainrefinement, achieved through processing by severe plastic deformation (SPD), maylead to improvements in themechanical properties of HEAs. Equal-channel angular pressing (ECAP) 10 and high-pressuretorsion (HPT) 11 arewell-established SPD proceduresfor achieving ultrafine and evennanostructured grains in metals and alloys but HPT is generally advantageous because, bycomparison with ECAP, it produces smaller grain sizes 12,13 and higherfractions of grain boundaries having high angles of misorientation 14.

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There are now several reports describing theinfluence of HPT processing on HEAs 15-24 and a single report documenting the processing ofan HEA by ECAP 25.  Furthermore, itwas shown earlier that processing aCoCrFeNiMn high-entropy alloy by HPT leads to significant hardening and grainrefinement and, in addition, post-deformation annealing (PDA) of thenanocrystalline CoCrFeNiMn HEA provides an excellent combination of highstrength and good ductility 19. Experiments also showed that ananocrystalline CoCrFeNiMn HEA processed by HPT exhibited excellent ductilityat elevated temperatures including superplastic elongations with a maximumelongation of >600% at a testing temperature of 973 K 22,26. One of the potential methods forincreasing the strength of HEAs with an fcc crystal lattice is doping bycarbon for the formation of high-strength fine carbides. Nevertheless, theinfluence of these carbides on the structure and properties of HEAs has receivedonly limited attention.  Accordingly, thepresent research was initiated to evaluate the effect of grain refinement dueto HPT on the microstructures of two CrFe2NiMnV0.

25 HEAs doped by carbon. Experimental materials and proceduresCrFe2NiMnV0.25C0.075and CrFe2NiMnV0.

25C0.125 high entropy alloys were prepared using a non-consumable vacuum arcmelting technique in a water-cooled copper crucible. The purities of thealloying elements were above 99.9 at.

%. After several remeltings (5 times) forhomogenization, the ingots were hot forged and then homogenized at 1323 K for 1h.  Then the samples were cold-rolled to~30 % thickness reduction followed by annealing at 1373 K for 1 h. To preventoxidation, all samples were sealed in vacuum quartz tubes filled with titaniumchips before the heat treatments.  Polisheddisks with diameters of 10 mm and thicknesses of ~0.8 mm were prepared from theannealed 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-constrainedconditions in which there is a small outflow of material around the peripheriesof the disks during the torsional straining 27.

After processing, each HPT disk was polished to amirror-like quality and hardness measurements were taken using a Vickersmicrohardness tester with a load of 500 gf and dwell times of 10 s. The averagemicrohardness values, Hv, were measured along randomly selected diameters oneach disk. These measurements were taken at intervals of ~0.5 mm and at everypoint the local value of Hv was obtained from the average of four separatehardness values. The phase constituents were determined using X-ray diffraction(XRD) employing Cu K? radiation (wavelength ? = 0.

154 nm) at 45 kV and a tubecurrent of 200 mA with Rigaku SmartLab equipment. The XRD measurements wereperformed over 2? angular ranges from 30° to 100° using a scanning step of0.01° and a scanning speed of 2º min-1.  Microstructural characterizations were carried outusing optical microscopy (OM) and transmission electron microscopy (TEM). Foilsfor TEM were prepared before PDA and after PDA at 823 K for 60 min using afocused ion beam (FIB) Zeiss Nvision 40 FIB facility at 3 mm from the CrFe2NiMnV0.25C0.125HEA disk centres in the normal sections of the disks sothat the normals of the images lay in the shear direction.

The TEM micrographswere obtained using a JEOL JEM-3010 microscope operating under an acceleratingvoltage of 300 kV. Experimental resultsInitialmicrostructure of  HEA Figure 1 shows the microstructure of theas-cast CrFe2NiMnV0.25C0.125HEA. The observations reveal an extended dendritic structure in Fig. 1(a)including darker dendrites and lighter interdendritic areas as shown in thehigher magnification image in Fig. 1(b).

The inter-dendrite spacing was aboutseveral micrometers and the average grain size was measured as ~300 mm. Themicrostructure contained approximately 6% of carbides (Cr23C6or Cr22.23Fe0.77C6) in the form of particles.

The microstructure of the CrFe2NiMnV0.25C0.075 HEA revealedsimilar regularities. It was shown earlier for the same alloy that the second phase in the interdendritic eutectic is chromium carbide wherein some of thechromium atoms are substituted by other elements 28.Hardnessevaluations before and after HPT processingFigure2 shows the results for the Vickers microhardnessmeasurements after processing through 5 turns with the average values of Hvplotted along each disk diameter and with the lower dashed lines at Hv ? 170and 180 corresponding to the initial hardness in the coarse- grained CrFe2NiMnV0.25C0.075and CrFe2NiMnV0.25C0.

125HEAs, respectively. The results clearly show thatthe hardness at the edge of the disk increases significantly after 5 turns by afactor of ~2.5 to Hv ? 430 and ~435 , with reference to the annealed condition,for the CrFe2NiMnV0.25C0.075and CrFe2NiMnV0.

25C0.125HEAs, respectively. These resultsdemonstrate that after 5 turns it is not possible to produce afully-homogeneous hardness distribution and instead the results show there is avery small area, within a radius of r < 1 mm, at the centres of thedisks where the hardness values are significantly lower (Hv ? 350) in both HEAs.  The results demonstrate also there is nosignificant difference between the hardness of samples processed by HPT through5 turns at radii larger than ~1 mm. Microstructuresbefore and after HPT processingXRDpatterns for bothinitial annealed samples not processed by HPT, designated as N = 0, andnear the edges of disks after HPT through totals of N = 5 turns for CrFe2NiMnV0.25C0.

075and CrFe2NiMnV0.25C0.125 are shown in Fig.

3(a) and Fig. 3(b), respectively. Inspection of patterns reveal that theannealed microstructures consist of an fcc phase marked with solidcircles and peaks marked with solid inverted triangles with the main peak at 2? ? 43.9º relatedto Cr23C6or the Cr22.23Fe0.77C6 phase.  The volume fraction of the secondary phase of chromium carbide in the initialcoarse-grained state was ~7%.

The patterns reveal also that HPT processing causes peak broadening inthe HEAs and it is important to note there is no evidence for the creation ofnew phases or the occurrence of any phase transformations during the HPTprocessing.The microstructure of the CrFe2NiMnV0.25C0.

125HEA is shown in Fig. 4 after 5 turns of HPT at a position near the edge areaswithout any particles in Fig. 4(a) and with a chromium carbide particle in Fig.

4(b). Inspection showed that the processed microstructure consisted of an arrayof equiaxed nanostructured grains having an average size of ~30 nm and withmany of the grains surrounded by diffuse or ill-defined grain boundaries as inFig. 4(a). In addition, strain contrast was visible in many of these nano-scalegrains and this was associated with the presence of dislocations.

Theappearance of this microstructure is typical of materials prepared using SPDtechniques and it is consistent with the presence of a large volume ofhigh-energy non-equilibrium boundaries 29. The arrangement of the diffraction spotsin semi-continuous circles in the selected area electron diffraction (SAED)pattern shown as an inset in Fig. 4(a) confirmed that the microstructurecontained an fcc phase in this area including boundaries having highangles of misorientation. A particle surrounded by matrix grains is shown inFig. 4(b) and energy dispersive X-ray spectroscopy (EDS) of the two regionsmarked A and B confirmed from the chemical compositions that these neighbouringareas were different phases.

Specifically, the chemical composition of region Acorresponded to the composition of the HEA matrix and the composition of regionB corresponded to chromium carbide which is consistent with the XRD results. Acomparison between the particle sizes before HPT (several microns) and afterHPT processing (<1 ?m) confirmed that the particles were crushed due to thehigh hydrostatic stresses imposed in the HPT processing. Microstructures after post-deformationannealingFigure 5 uses XRD to demonstrate the microstructural evolution in theCrFe2NiMnV0.

25C0.125 samples after annealingat 823 and 1073 K for 60 min.  The effectof this subsequent annealing on the initial annealed microstructure is shown inFig. 5(a) and the effect on the HPT-processed material is shown in Fig.5(b).  Inspection reveals the appearanceof some additional peaks of very low intensities marked with open invertedtriangles in the samples annealed at 823 K in both the initial annealed and theHPT-processed samples.

  The crystalstructure of this new phase was identified as tetragonal with latticeparameters of a ? 8.8 Å and c ? 4.5 Å where this corresponds tothe ?-phase which is a hard Cr-rich phase reported earlier in some HEAsprocessed by SPD 16-22,24.

  Theseresults indicate that the volume fraction of the ?-phase was ~3% and ~6%in the annealed and HPT-processed samples, respectively, and close inspectionshows an absence of this phase after annealing at 1273 K.  A scanning transmission electronmicroscopy (STEM) image is shown in Fig. 6 for a CrFe2NiMnV0.25C0.125sample after HPT for 5 turns followed by PDA at 823 K for 60 min. Thus,in this condition the microstructure contains essentially equiaxed grains withan estimated mean grain size of ~200 nm. The STEM micrograph shows also the existence of a precipitated phasewhich is darker in appearance and, according to XRD results related to thisannealing temperature, appear to correspond to precipitates of a Cr-richphase.  It is apparent also that thisCr-rich phase is distributed reasonably homogenously throughout themicrostructure in Fig.

6 with an average size of ~100 nm.Table 1 records the measured values of the microhardness for the initialcoarse-grained CrFe2NiMnV0.25C0.125 HEA andfor the HPT-processed samples both in their initial conditions and afterannealing at temperatures of 823 and 1273 K for a period of 60 min. The resultsdemonstrate that the hardness of the nanocrystalline HEA increases slightly upto Hv ? 555 after annealing at 823 K and then decreases rapidly with increasingannealing temperatures up to 1273 K.  Atthis latter temperature, the hardness is ~195 which is close to the value of Hv» 180 for the initial coarse-grained condition.

  It is interesting to note also that theannealed sample tends to show a similar trend but with a maximum hardness ofonly Hv ? 190 after annealing at 823 K. This behavior suggests that the precipitates form at temperatures up to823 K in both conditions and this is consistent with the XRD results as shownin Fig. 5. DiscussionEffect of HPT and carbon content on the microstructures of HEAs The XRD and hardness results show that there is no significantdifference between the CrFe2NiMnV0.

25C0.075 andCrFe2NiMnV0.25C0.125 HEAs after HPTprocessing.  The hardness of CrFe2NiMnV0.

25C0.125is marginally larger than CrFe2NiMnV0.25C0.075with a lower carbon content but the volume fractions of chromium carbide arealmost the same at ~6% in both HEAs. The microstructural observations of theCrFe2NiMnV0.25C0.125 HEA after HPT processingreveal exceptional grain refinement to ~30 nm with a dispersion of chromiumcarbide nanoparticles within the fcc matrix.

  It has been shown that grain refinement inthe single-phase HPT-processed CoCrFeNiMn alloy is more efficient with grainsizes of ~10 nm 19 by comparison with the CrFe2NiMnV0.25C0.125HEA.  Thus, the presence ofcarbides in the CrFe2NiMnV0.25C0.125 HEAproduces barriers that appear to restrict further grain refinement 30,31.

 Microstructuralevolution of the fcc HEA was studied earlier 5,32 and it was foundthat the most significant plasticity mechanism at low homologous temperatureswas gliding of full<110> dislocations on the primary {111} planes with part of the<110> dislocations splitting into 1/6<112> Shockley partialscontaining stacking faults.  Hardprecipitates in the fcc HEAs, such as particles of the ?-phase,act as strong barriers for dislocation motion so that most plasticity takesplace in the matrix fcc phase 33. The mechanisms of plasticity in fccnanocrystalline materials are not fully understood at the present time butgenerally it is considered that grain boundary sliding is especially importantat grain sizes below ~10-15 nm whereas for larger grains the motion of partialdislocations becomes dominant and this changes to the gliding of fulldislocations when the grain size increases above ~100 nm 34.  Below ~100-200 nm the plasticity iscontrolled by the nucleation and annihilation of dislocations at grainboundaries.  At present no specificinformation is available on the plasticity mechanisms in nanocrystalline fccHEAs but it is suggested that probably their behaviour is similar to themechanisms in conventional fcc nanocrystalline alloys. Accordingly, forHPT-processed nanocrystalline CrFe2NiMnV0.25C0.075and CrFe2NiMnV0.

25C0.125 HEAs with averagegrain sizes of ~25 nm it is reasonable to anticipate that the splitting ofdislocations into partials plays an important role by comparison with theircoarse-grained counterparts. If the strength of the nanocrystalline materials iscontrolled by the nucleation of dislocations at grain boundaries which dependsdirectly on the grain size, then the other contributions to yield strengthtypical for coarse-grained materials, such as hard particles of second phasesand interstitial impurities 35, should be less important in thenanocrystalline state. Such behaviour can be observed by comparing thehardening by carbides in the CrFe2NiMnV0.25C0.075and CrFe2NiMnV0.

25C0.125 HEAs in Fig. 2 withearlier reports for the single phase CoCrFeNiMn HEA 16,19.  In the coarse-grained state the hardening bycarbides leads to higher strength and higher hardness for CrFe2NiMnV0.25C0.125by comparison with CoCrFeNiMn while in the nanocrystalline state the singlephase CoCrFeNiMn HEA has a higher hardness of ~450 Hv due to the smaller grainsize of ~10 nm 19 compared with an average grain size of ~30 nm and ahardness of ~435 Hv in the CrFe2NiMnV0.25C0.125HEA.

  It is concluded, therefore, thathardening by carbides in the CrFe2NiMnV0.25C0.125HEA is less effective than hardening through the smaller grain size in theCoCrFeNiMn HEA.Thermal stability of HEAs after PDAThe microstructures of the two HEAs consist of an fccphase and chromium carbide particles before and after HPT processing butfurther annealing at 823 K leads to formation of new precipitates.

  This isconsistent with the substantial increase in hardness upon annealing and withthe XRD results.  Formation of precipitates in a single phaseHEA after annealing within special temperature ranges is a well-knownphenomenon 36-39. For example, it was shown that the CoCrFeMnNi alloy has asingle-phase fcc structure above 873 K but a mixture of fcc and bccphases, or under some conditions a ? phase with a tetragonal crystalstructure, below 873 K 19.  Inaddition, CoCrFeNiMnCx (x = 0.1, 0.

175, 0.25) HEAs demonstrate asignificant increase in hardness after annealing in the temperature range of875-1275 K due to the formation of precipitates 31. The present resultsconfirm the formation of a multi-phase nanostructured HEA after PDA at 823 Kconsisting of precipitates distributed within the microstructure. Closeinspection of XRD results (Fig. 5) reveals the ? phase in the initial annealed and HPT-processedsamples after annealing at 823 K. Nevertheless, these results indicate that the volume fraction ofprecipitates in the HPT-processed sample is higher than in the initial annealedsample. Basically, it iswell known that HEAs have sluggish diffusion which affects diffusion controlledmechanisms such as precipitating and grain coarsening 40-42. It appears that thelarge number of grain boundaries and imposed defects in the nanocrystallineHPT-processed promote fast diffusion pathways and also as preferentialnucleation sites for the formation of precipitates.

Thus, the severeplastic deformation leads more quickly to the formation of stable precipitatescompared with fullyannealed samples. Inspection of the hardness results in Table 1 showsthe hardness decreases significantly above 773 K and up to 1273 K and this is dueto 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 annealedcondition due to graincoarsening. Therefore, the dissolution of the precipitates plays an importantrole in the stability of the microstructure and grain coarsening during annealing at 823K.