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Article Effects of on Microstructure and Mechanical Properties of Metastable Powder CoCrFeNiMo0.2 High Entropy

Cui Zhang 1, Bin Liu 1,*, Yong Liu 1, Qihong Fang 2, Wenmin Guo 1,* and Hu Yang 3

1 State Key Laboratory for Powder Metallurgy, Central South University, Changsha 410083, China; [email protected] (C.Z.); [email protected] (Y.L.) 2 College of Mechanical and Vehicle Engineering, Hunan University, Changsha 410082, China; [email protected] 3 Yuanmeng Precision Technology (Shenzhen) Institute, Shenzhen 518055, China; [email protected] * Correspondence: [email protected] (B.L.); [email protected] (W.G.); Tel.: +86-731-8887-7669 (B.L.); +86-739-530-5016 (W.G.)  Received: 15 February 2019; Accepted: 24 April 2019; Published: 30 April 2019 

Abstract: A CoCrFeNiMo0.2 high entropy alloy (HEA) was prepared through powder metallurgy (P/M) process. The effects of annealing on microstructural evolution and mechanical properties of P/M HEAs were investigated. The results show that the P/M HEA exhibit a metastable FCC single-phase structure. Subsequently, annealing causes precipitation in the grains and at the grain boundaries simultaneously. As the increases, the size of the precipitates grows, while the content of the precipitates tends to increase gradually first, and then decrease as the annealing temperature goes up to 1000 ◦C. As the annealing time is prolonged, the size and content of the precipitates gradually increases, eventually reaching a saturated stable value. The mechanical properties of the annealed alloys have a significant correspondence with the precipitation behavior. The larger the volume fraction and the size of the precipitates, the higher the strength and the lower the of the HEA. The CoCrFeNiMo0.2 high entropy alloy, which annealed at 800 ◦C for 72 h, exhibited the most excellent mechanical properties with the ultimate tensile strength of about 850 MPa and an elongation of about 30%. Nearly all of the annealed HEAs exhibit good strength– combinations due to the significant precipitation enhancement and nanotwinning. The separation of the coarse precipitation phase and the matrix during the deformation process is the main reason for the formation of micropores. Formation of large volume fraction of micropores results in a decrease in the plasticity of the alloy.

Keywords: powder metallurgy; high entropy alloy; microstructure; precipitation strengthening; mechanical properties

1. Introduction Designing strong and ductile metals has been among the most ambitious goals for metallurgists [1–4]. During the past decade, a new concept of (HEAs), which is broadly defined based on the high entropy effect of alloys with multi-components, has attracted great attention due to their unique superior properties, such as high strength, good ductility, high thermal stability, and oxidation resistance [5–7]. The most studied HEAs can be divided into two categories: one is FCC HEAs, with Fe, Co, Ni, Cr, Mn and other transition elements as main components, and the other is BCC HEAs, with refractory metals as main components [8–10]. Previous studies have demonstrated that FCC structured HEAs generally exhibited high tensile elongation but relatively low

Entropy 2019, 21, 448; doi:10.3390/e21050448 www.mdpi.com/journal/entropy Entropy 2019, 21, 448 2 of 9 strength, while BCC alloys are the opposite, although some relatively ductile BCC refractory HEAs have been reported [11–13]. Simultaneous improvement of strength and ductility of alloys has been very challenging. Although single-phase FCC high-entropy alloys struggle to meet the high strength requirements of engineering materials, such alloys exhibit high work rates and uniform deformation behavior [14,15]. These characteristics make high-entropy alloys prone to be an excellent composite matrix. Yang et al. [1] reported that the FCC based (CoFeNi)86-Al7Ti7 alloy can be significantly strengthened by the L12 multicomponent intermetallic nanoparticles and exhibit superior strength of 1.5 GPa and ductility as high as 50% in tension at ambient temperature. Liu et al. [16] reported that the FCC CoCrFeNiNbx can be significantly strengthened by the Nb-enriched Laves phase with the HCP structure. He et al. [17] reported that the addition of Al to the CoCrFeNiMn alloy can form a BCC phase, resulted in a high tensile strength up to 1174 MPa. The above studies show that the formation of high reinforcing particles is an effective way to prepare high performance high entropy composite alloys. Presently, different kinds of topologically close-packed (TCP) phases, such as σ, µ, Laves, etc. are observed in HEAs [18,19]. Adding this reinforcing phase to the FCC high entropy alloys can significantly increase their strength, but at the same time significantly reduce their plasticity [20]. Based on the precipitation strengthening effect, the most important thing is to control the morphology and distribution of the precipitated phase, and then to alleviate their harmful effects on ductility. As reported, the size and distribution of the reinforcing phase is very sensitive to heat treatment conditions. Gwalani et al. [21] reported that the L12 Ni3 (Ti, Al) nano-precipitates in the Al0.3CrFeCoNi alloy can only be stably present in the temperature range of 500–600 ◦C. As the above ~700 ◦C, these precipitates are dissolved and replaced by coarser ordered B2 precipitates. Liu et al. [14] reported that the precipitated nano σ phase formed during the heat treatment at 850–900 ◦C can simultaneously increase the strength and plasticity of the CoCrFeNiMo0.3 alloy. Ming et al. [22] reported that the Cr15Fe20Co35Ni20Mo10 HEA exhibit a superb strength–ductility combination by of nanoscale precipitates of Mo-rich µ phase. It is suggested that a reasonable heat treatment process is essential for the formation of nano-precipitates, generally resulting in the improved comprehensive mechanical properties of alloys. Powder metallurgy technology is one of the effective methods to avoid component segregation and obtain high performance composite materials. There have been many reports on powder metallurgy high entropy alloys [23–25]. Due to the higher cooling rate during the preparation process, these alloys typically exhibit an equiaxed microstructure. Presently, the effects of precipitation strengthening on the mechanical properties of CoCrFeNi based HEAs prepared by powder metallurgy have not yet been reported. In the present research, a CoCrFeNiMo0.2 high entropy alloy was prepared by P/M process. A systematic study on precipitation behavior of the reinforcing phase is obtained. The effects of annealing on microstructure and mechanical properties of powder metallurgy CoCrFeNiMo0.2 high entropy alloy were also discussed.

2. Experimental Procedures

The CoCrFeNiMo0.2 spherical powders were obtained by an atomization process with high purity Fe Co Cr Ni and Mo raw materials. These raw materials were firstly melted in a . And then, the melt dropped through a tube and was atomized in high purity Ar with an atomization pressure was 4 MPa. The chemical composition and oxygen content of the atomized powders was characterized by chemical methods and the fusion method on a Leco O/N analyzer (LECO TCH 600) respectively. Subsequently, the CoCrFeNiMo0.2 alloy was prepared by a hot extrusion process with the atomized powders. The dimensions of the stainless- mold used in the hot extrusion process is d60 150 mm3. × The powder is first loaded into a can, pre-heated at 1473 K for 60 min, and sealed under Entropy 2019, 21, 448 3 of 9 vacuum.Entropy 2019 The, 21, x enclosed FOR PEER powders REVIEW were immediately subjected to hot extrusion with an extrusion3 ratio of 9 of 6 and a velocity of ~10 mm/s on a 2500 T hydraulic press. The as-extruded alloy was annealed at was annealed at different temperatures range from in the range between 700 and to 1000 °C in different temperatures range from in the range between 700 and to 1000 ◦C in vacuum for different vacuumtimes range for different from 2 to times 72 h respectively, range from 2 and to 72 then h respectively, water quenched. and then water quenched. 3 Tensile samples withwith dd44 × 15 mm mm3 werewere prepared prepared with with the the as-extruded as-extruded and annealed HEA alloys × along thethe extrusionextrusion direction direction (ED). (ED). Microstructures Microstructure ofs these of these alloys alloys were analyzedwere analyzed using a fieldusing emission a field emissionscanning scanning electron microscopeelectron microscope (FESEM) (FESEM) (FEI Nova (F Nano-230,EI Nova Nano-230, Hillsboro, Hillsboro, OR, USA) OR, equipped USA) equipped with an withelectron an backscatteredelectron backscattered diffraction systemdiffraction (EBSD). system The size(EBSD). and volumeThe size fraction and ofvolume the precipitated fraction of phase the wereprecipitated obtained phase by image were analysis obtained using by image Imagepro analysis software. usingPhase Imagepro structures software. were Phase identified structures by an X-raywere identifieddiffractometer by an (XRD) X-ray (Rigaku diffractometer D/MAX-2250, (XRD) Tokyo, (Rigaku Japan) D/MAX-2250, with a Cu/Ka Tokyo, radiation. Japan) Tensile with tests a Cu/Ka were radiation. Tensile tests were performed with3 a loading strain rate of 10−3/s on an Instron 3369 testing performed with a loading strain rate of 10− /s on an Instron 3369 testing machine at room temperature. Themachine standard at room bright-field temperature. images andThe distandardffraction patternsbright-field were images obtained and using diffraction a transmission patterns electron were obtainedmicroscope using (Tecnai a transmission G2 F20 S-TWIN, electron FEI, microscope Hillsboro, (Tecnai OR, USA). G2 F20 S-TWIN, FEI, Hillsboro, OR, USA).

3. Results Results and and Discussion

3.1. Microstructures Microstructures

Figure1 a1a shows shows the the inverse inverse pole pole figure figure (IPF) (IPF) of the of as-extruded the as-extruded CoCrFeNiMo CoCrFeNiMo0.2 HEA.0.2It HEA. is obvious It is obviousthat the extrudedthat the alloyextruded exhibits alloy an exhibits equiaxed an grain equiaxed structure grain with structure an average with grain an average size of approximately grain size of approximately20 µm. The relative 20 μm. density The relative of the as-extrudeddensity of the alloy as-extruded is approximately alloy is approximately 99.5%. The consolidated 99.5%. The consolidatedmicrostructure microstructure is fully recrystallized, is fully recrystallized, indicating that indicating recrystallization that recrystallization occurred during occurred and after during the andextrusion. after the Figure extrusion.1b shows Figure the XRD 1b shows patterns the of XR theD Ppatterns/M CoCrFeNiMo0.2 of the P/M CoCrFeNiMo0.2 HEA, where the HEA, alloy where shows theclearly alloy a singleshowsFCC clearly structure. a single FCC structure.

Figure 1. (a) IPF mapmap ofof thethe PP/M/M CoCrFeNiMoCoCrFeNiMo0.20.2 high entropy alloy (HEA), ( b)) XRD patterns of the P/MP/M CoCrFeNiMo0.2 alloy.

Figure2 2 demonstrate demonstrate thethe microstructuresmicrostructures ofof thethe CoCo CrCr FeFe NiNi MoMo 0.20.2 HEAsHEAs annealedannealed atat didifferentfferent temperatures in the range of 700–1000 ◦°CC for 72 h respectively.respectively. These annealed alloys are mainly composed of distinct grey matrix, black pores and white areas. It is apparent from Figure 2 2aa toto dd thatthat the size ofof thesethese whitewhite areasareas gradually gradually increases increases as as the the annealing annealing temperature temperature increased. increased. As As shown shown in µ inFigure Figure2b, 2b, white white areas areas generally generally appear appear at the at grainthe grain boundaries boundaries and and the sizethe size is less is less than than 1 m. 1 μ Asm. theAs theannealing annealing temperature temperature increased increased up toup1000 to 1000◦C, °C, the the size size of theseof these white white areas areas is rapidlyis rapidly coarsened coarsened to to3–5 3–5 microns. microns. In In addition, addition, the the volume volume fraction fraction of theseof these white white areas areas tends tends to increase to increase gradually gradually first first and andthen then decrease, decrease, as the as annealing the annealing temperature temperature goes goes up to up 1000 to 1000◦C. °C. In order to further identify the structure of the P/M CoCrFeNiMo0.2 HEA, we performed the TEM analysis on the precipitated phase. It is reported that the σ phase and µ phase in the CoCrFeNiMox alloy systems is corresponding to the stoichiometric (Cr,Mo)(Co,Fe,Ni) and (Mo,Cr)7(Co,Fe,Ni)6 respectively [26,27]. The EDS analysis results in Table1 clearly indicates that the chemical composition of the precipitates contains a high concentration of Mo, which is very close to the σ phase reported by Shun et al. [27,28]. The selected electron diffraction pattern embedded in the upper right corner of

Entropy 2019, 21, 448 4 of 9

EntropyFigure 20193 also, 21, confirms x FOR PEER that REVIEW the white precipitates are σ phase. This result is also consistent with4 of the 9 Entropycalculated 2019, pseudo21, x FOR binaryPEER REVIEW (CoCrFeNi)-Mo [14]. 4 of 9

Figure 2. SEM images ofof thethe PP/M/M CoCrFeNiMo CoCrFeNiMo0.20.2HEAs HEAs annealed annealed at at di differentfferent temperatures temperatures for for 72 h.72 (h.a) Figure 2. SEM images of the P/M CoCrFeNiMo0.2 HEAs annealed at different temperatures for 72 h. (700a) 700◦C, °C, (b) ( 800b) 800◦C, °C, (c) ( 900c) 900◦C, °C, and and (d) ( 1000d) 1000◦C. °C. (a) 700 °C, (b) 800 °C, (c) 900 °C, and (d) 1000 °C.

In order to further identify the of the P/M CoCrFeNiMo0.2 HEA, we performed InTable order 1. EDS to further analysis identify results forthe the crystal P/MCoCrFeNiMo structure of theHEA P/M annealed CoCrFeNiMo at 700 0.2C HEA, for 72 hwe (Spots performed in the TEM analysis on the precipitated phase. It is reported0.2 that the σ phase◦ and μ phase in the the TEMFigure analysis3). on the precipitated phase. It is reported that the σ phase and μ phase in the CoCrFeNiMox alloy systems is corresponding to the stoichiometric (Cr,Mo)(Co,Fe,Ni) and CoCrFeNiMox alloy systems is corresponding to the stoichiometric (Cr,Mo)(Co,Fe,Ni) and (Mo,Cr)7(Co,Fe,Ni)6 respectively [26,27]. The ChemicalEDS analysis Composition results in (at.%) Table 1 clearly indicates that the (Mo,Cr)7(Co,Fe,Ni)Location6 respectively [26,27]. The EDS analysis results in Table 1 clearly indicates that the chemical composition of the precipitates contains a high concentration of Mo, which is very close to chemical composition of the precipitatesMo L contains Cr K a high Feconcentration K Co Kof Mo, which Ni K is very close to the σ phase reported by Shun et al. [27,28]. The selected electron diffraction pattern embedded in the the σ phase reportedEDS spot by Shun 1 et al. 34.56 [27,28]. The 18.03 selected electron 20.38 diffraction 17.53 pattern 9.49 embedded in the upper right corner of Figure 3 also confirms that the white precipitates are σ phase. This result is also upper right cornerEDS of spot Figure 2 3 also 6.18 confirms that 22.08 the white 25.71 precipitates 24.51 are σ phase. 21.52 This result is also consistent with the calculated pseudo binary (CoCrFeNi)-Mo phase diagram [14]. consistent with the calculated pseudo binary (CoCrFeNi)-Mo phase diagram [14].

Figure 3. TEM image of the P/M CoCrFeNiMo0.2 HEA annealed at 700 ◦C for 72 h. Figure 3. TEM image of the P/M CoCrFeNiMo0.2 HEA annealed at 700 °C for 72 h. Figure 3. TEM image of the P/M CoCrFeNiMo0.2 HEA annealed at 700 °C for 72 h. The microstructures of the CoCrFeNiMo0.2 HEAs annealed at 800 ◦C for various times are Table 1. EDS analysis results for the P/M CoCrFeNiMo0.2 HEA annealed at 700 °C for 72 h (Spots in illustratedTable in1. EDS Figure analysis4. It is results apparent for the that P/M prolongation CoCrFeNiMo of0.2 the HEA annealing annealed time at 700 promotes °C for 72 the h precipitation(Spots in Figure 3). of theFigureσ phase. 3). As shown in Figure4b, σ phase began to appear at the of the matrix as the annealing time was 4 h. σ phase grows withChemical annealing Composition time and (at.%) changes its morphology. Location Chemical Composition (at.%) Location Mo L Cr K Fe K Co K Ni K Mo L Cr K Fe K Co K Ni K EDS spot 1 34.56 18.03 20.38 17.53 9.49 EDS spot 1 34.56 18.03 20.38 17.53 9.49 EDS spot 2 6.18 22.08 25.71 24.51 21.52 EDS spot 2 6.18 22.08 25.71 24.51 21.52

Entropy 2019,, 21,, xx FORFOR PEERPEER REVIEWREVIEW 5 of 9

The microstructures of the CoCrFeNiMo0.20.2 HEAsHEAs annealedannealed atat 800800 °C°C forfor variousvarious timestimes areare illustratedillustrated inin FigureFigure 4.4. ItIt isis apparentapparent thatthat prolongation of the annealing time promotes the precipitation of the σ phase.phase. AsAs shownshown inin FigureFigure 4b,4b, σ phasephase beganbegan toto appearappear atat thethe graingrain boundaryboundary Entropyof the 2019matrix, 21, 448as the annealing time was 4 h. σ phasephase growsgrows withwith annealingannealing timetime andand changeschanges5 ofitsits 9 morphology.

Figure 4. SEMSEM images of the P/M P/M CoCrFeNiMo 0.20.2 HEAsHEAsHEAs annealedannealed atat 800800 °C°C◦C forfor differentdifferent different times: times: (( aa)) 2 2 h,h, h, ((b)) 44 h,h, ((cc)) 88 h,h, ((dd)) 161616 h,h,h, (((ee))) 484848 h,h,h, andandand (((ff))) 727272 h.h.h.

Figure5 5 shows shows a statisticala statistical analysis analysis on on the the variation variation of the of sizethe andsize volumeand volume fraction fraction of the ofσ phasethe σ withphase the with annealing the annealing temperature temperature and time. and Astime. shown As shown in Figure in Figure5a, the 5a, size the of size the ofσ thephase σ phasephase is less isis than lessless than0.5 µ m0.5 as μ them as annealing the annealing temperature temperature at 700 at◦C. 700 As °C. the As temperature the temperature increases, increases, the size the of size the ofσ phase the σ phaserapidly rapidly grows grows to the micronto the micron level. Whenlevel. theWhen annealing the annealing temperature temperature increases increases to 1000 to◦C, 1000 the °C, size the of sizethe σ ofphase the σ reachesphasephase reachesreaches to 3.7 µ totom. 3.73.7 As μm. seen As from seen Figurefrom Figure5b, the 5b, content the co ofntent the of precipitated the precipitated phase phase is also is alsoclosely closely related related to the to annealing the annealing time. time. As the As annealingthe annealing time time is prolonged, is prolonged, the volumethe volume fraction fraction of the of precipitatesthe precipitates gradually gradually increases, increases, eventually eventually reaching reaching a saturated a saturated stable value. stable Therefore, value. obtainingTherefore, a obtaininguniformly a dispersed uniformly nanoprecipitate dispersed nanoprecipitate phase in the CoCrFeNiMo phase in the0.2 HEACoCrFeNiMo is strictly0.20.2 controlledHEAHEA isis strictlystrictly by the controlledheat treatment by the process. heat treatment process.

Figure 5.5. (((aa))) VariationVariation Variation ofof of thethe averageaverage sizesize andand volume volume fraction fraction of of σ precipitateprecipitate withwith annealingannealing temperatures;temperatures; ((b)) VariationVariation ofofof thethethe volumevolumevolume fractionfractionfraction of ofof σ σ precipitate precipitateprecipitate with withwith annealing annealingannealingtimes timestimesat atat800 800800◦ °C.°C.C. 3.2. Mechanical Properties 3.2. Mechanical Properties Figure6a shows the engineering stress-strain curves of the as-extruded and annealed Figure 6a shows the engineering stress-strain curvescurves ofof thethe as-extrudedas-extruded andand annealedannealed CoCrFeNiMo0.2 HEA at different temperatures for 72 h. The evolution of tensile strength and plasticity CoCrFeNiMo0.20.2 HEAHEA atat differentdifferent temperaturestemperatures forfor 7272 h.h. TheThe evolutionevolution ofof tensiletensile strengthstrength andand with annealing temperature was also statistically calculated in Figure6b. A typical elasto-plastic plasticity with annealing temperature was also statistically calculated in Figure 6b. A typical elasto- deformation behavior is remarked. All these annealed alloys exhibited a long stage. plastic deformation behavior is remarked. All these annealed alloys exhibited a long work hardening The yield strength, ultimate tensile strength and elongation to failure of the as-extruded CoCrFeNiMo stage. The yield strength, ultimate tensile strength and elongation to failure of the as-extruded0.2 HEA were about 400 MPa, 781 MPa and 55.6% respectively. As the annealing temperature increases from 700 ◦C to 900 ◦C, it is apparent that the yield strength and ultimate tensile strength of these annealed alloys were gradually improved. However, the ductility of these alloys has also dropped significantly. Since the annealing temperature increases up to 1000 ◦C, the yield strength and ultimate Entropy 2019, 21, x FOR PEER REVIEW 6 of 9

EntropyCoCrFeNiMo 2019, 21, x0.2 FOR HEA PEER were REVIEW about 400 MPa, 781 MPa and 55.6% respectively. As the annealing6 of 9 temperature increases from 700 °C to 900 °C, it is apparent that the yield strength and ultimate tensile CoCrFeNiMostrength of these0.2 HEA annealed were alloysabout were400 MPa,gradually 781 MPaimproved. and 55.6%However, respectively. the ductility As ofthe these annealing alloys temperature increases from 700 °C to 900 °C, it is apparent that the yield strength and ultimate tensile Entropyhas also2019 dropped, 21, 448 significantly. Since the annealing temperature increases up to 1000 °C, the yield6 of 9 strengthstrength ofand these ultimate annealed tensile alloys strength were ofgradually the alloy improved. is significantly However, decreased, the ductility while theof theseplasticity alloys is hascorrespondingly also dropped increased significantly. to as Since high theas approximatelyannealing temperature 65%. Although increases such up a totradeoff 1000 °C, relationship the yield strengthtensilebetween strength andthe ultimatestrength of the alloy tensileand isplasticity significantlystrength is of convention the decreased, alloy isal, while significantlythe comprehensive the plasticity decreased, is correspondinglymechanical while the properties plasticity increased ofis correspondingly increased to as high as approximately 65%. Although such a tradeoff relationship tothese as highannealed as approximately CoCrFeNiMo 65%.0.2 alloys Although are significantly such a tradeo betterff thanrelationship those of betweenthe as-extruded the strength one. and betweenplasticity the is conventional, strength and the plasticity comprehensive is convention mechanicalal, the properties comprehensive of these mechanical annealed CoCrFeNiMo properties 0.2of thesealloys annealed are significantly CoCrFeNiMo better0.2 than alloys those are of significantly the as-extruded better one. than those of the as-extruded one.

Figure 6. (a) Room-temperature engineering stress-strain curves for these HEAs annealed at different temperatures; (b) variation tendency of ultimate tensile strength and elongation of the annealed alloys Figure 6. (a) Room-temperature engineering stress-strain curves for these HEAs annealed at different Figurewith at 6. different (a) Room-temperature annealing temperatures. engineering stress-strain curves for these HEAs annealed at different temperatures; (b) variation tendency of ultimate tensile strength and elongation of the annealed alloys temperatures; (b) variation tendency of ultimate tensile strength and elongation of the annealed alloys with at different annealing temperatures. withFigure at different7a shows annealing the engineeringtemperatures. stress-strain curves of the as-extruded and annealed CoCrFeNiMo0Figure7a.2 shows HEA at the 800 engineering °C with different stress-strain annealing curves times. ofThe the evolution as-extruded of tensile and strength annealed and Figure 7a shows the engineering stress-strain curves of the as-extruded and annealed CoCrFeNiMo0plasticity with .2annealingHEA at 800 time◦C was with also diff statisticallyerent annealing calculated times. in The Figure evolution 7b. With of tensilethe prolongation strength and of CoCrFeNiMo0.2 HEA at 800 °C with different annealing times. The evolution of tensile strength and plasticityannealing with time, annealing the yield time strength was alsoand statisticallyultimate tens calculatedile strength in Figureof these7b. annealed With the alloys prolongation gradually of plasticity with annealing time was also statistically calculated in Figure 7b. With the prolongation of annealingincrease, while time, the the plasticity yield strength decreases and gradually. ultimate tensile Combined strength with of the these Figure annealed 5, it is alloysobvious gradually that the annealing time, the yield strength and ultimate tensile strength of these annealed alloys gradually increase,evolutionary while trend the plasticity of the decreasesmechanical gradually. properties Combined of the with annealed the Figure alloys5, it ishas obvious a significant that the increase, while the plasticity decreases gradually. Combined with the Figure 5, it is obvious that the evolutionarycorrespondence trend with of thethe mechanical precipitate propertiesbehavior of of σ the phase. annealed The alloysenhanced has ayield significant strength correspondence and ultimate evolutionary trend of the mechanical properties of the annealed alloys has a significant withtensile the strength precipitate is mainly behavior attributed of σ phase. to the The high enhanced content yield and grain strength growth and ultimateof the σ tensilephase. strengthThe larger is correspondence with the precipitate behavior of σ phase. The enhanced yield strength and ultimate mainlythe volume attributed fraction to and the highthe size content of the and σ phase, the higher of the theσ strengthphase. The while larger the thelower volume the plasticity fraction tensile strength is mainly attributed to the high content and grain growth of the σ phase. The larger andof the the alloy. size of the σ phase, the higher the strength while the lower the plasticity of the alloy. the volume fraction and the size of the σ phase, the higher the strength while the lower the plasticity of the alloy.

Figure 7. (a) Room-temperature engineering stress-strain curvescurves forfor thethe HEAsHEAs annealedannealed atat 800800 ◦°CC for variousvarious times;times; ((bb)) variationvariation tendencytendency ofof ultimate tensiletensile strengthstrength and elongation of the annealed HEAs Figureat didifferentfferent 7. (a )annealingannealing Room-temperature times.times. engineering stress-strain curves for the HEAs annealed at 800 °C for various times; (b) variation tendency of ultimate tensile strength and elongation of the annealed HEAs Figure8 shows the fracture surface of CoCrFeNiMo 0.2 HEA annealed at different temperatures atFigure different 8 shows annealing the times.fracture surface of CoCrFeNiMo0.2 HEA annealed at different temperatures in the range from 700 C to 1000 C for 72 h. The presence of a large number of dimples on the in the range from 700 °C◦ to 1000 °C ◦for 72 h. The presence of a large number of dimples on the fracture fractureFigure surface 8 shows indicates the fracture that these surface alloys of have CoCrFeNiMo good deformability0.2 HEA annealed [29]. As theat different annealing temperatures temperature inincreases, the range it from can be 700 found °C to that 1000 the °C sizefor 72 of h. the The dimples presence becomes of a large significantly number of bigger.dimples Simultaneously, on the fracture obvious signs of broken phase of the precipitated phase can be found at the bottom of the dimple. It is found that the separation of the precipitation phase and the matrix during the deformation process is the main reason for the formation of micropores. Due to the difference between the elastic modulus, Entropy 2019, 21, x FOR PEER REVIEW 7 of 9 surface indicates that these alloys have good deformability [29]. As the annealing temperature increases, it can be found that the size of the dimples becomes significantly bigger. Simultaneously, obvious signs of broken phase of the precipitated phase can be found at the bottom of the dimple. It obviousEntropy 2019 signs, 21, 448of broken phase of the precipitated phase can be found at the bottom of the dimple.7 of It 9 is found that the separation of the precipitation phase and the matrix during the deformation process is the main reason for the formation of micropores. Due to the difference between the elastic modulus, stress concentration is likely to occur at the inte interfacerface between the precipitated phase and the matrix during the deformation process. The The size size of of the mi microporescropores increases with this coarse precipitated phase, indicating crack propagation in the localized region of the alloy was significantly significantly affected affected by these precipitatedprecipitated phase.phase. FormationFormation of of large large volume volume fraction fraction of of micropores micropores results results in ain decrease a decrease in the in theplasticity plasticity of the of alloy.the alloy.

Figure 8. The morphologies of the fractured surface of the PP/M/M CoCrFeNiMo0.2 HEAs annealed at differentdifferent temperatures for 7272 h.h. (a) 700 ◦°C,C, ( b) 800 °C,◦C, ( c)) 900 900 °C,◦C, and and ( (dd)) 1000 1000 °C.◦C.

Figure 99 illustratesillustrates thethe TEMTEM bright-fieldbright-field imagesimages ofof thethe annealedannealed CoCrFeNiMo0.2CoCrFeNiMo0.2 alloyalloy afterafter tensile deformation. It It can can be be seen seen that that a a large large nu numbermber of deformed twins formed during the tensile deformation process. The The interaction interaction of of the the precip precipitates,itates, deformed twins and grain boundary results in the pileuppileup ofof .dislocations. The The σσ precipitatesprecipitates did did not undergo plastic deformation during the slip, resulting in high strength and work hardening rate of the alloy. As can be seen from Figure 99b,b, thethe nanoprecipitatenanoprecipitate phasephase preferentiallypreferentially precipitatesprecipitates atat thethe graingrain boundary.boundary. InIn addition,addition, deformed twins also also play play an an important important role role in in im improvingproving the the strength strength and and toughness toughness of of the the alloy. alloy. The twin boundaries have a significantsignificant hindrancehindrance to thethe slipslip ofof dislocations.dislocations. Thus, these annealed CoCrFeNiMo0.20.2 HEAsHEAs plastically plastically deform deform via via dislocatio dislocationn gliding and nanotwinning,nanotwinning, significantly significantly enhancing strain hardening capability.

Figure 9. TEM bright-field images of the fractured CoCrFeNiMo0.2 HEA annealed at 700 ◦C for 72 h: (a) showing the nano twinning and (b) showing the interaction of dislocations with nanoscale σ precipitates. Entropy 2019, 21, 448 8 of 9

4. Conclusions

In the present research, a CoCrFeNiMo0.2 HEA is prepared using P/M method. The P/M HEA was annealed at different temperatures (700–1000 ◦C) for different times (2–72 h). The following conclusions are drawn: 1. The P/M CoCrFeNiMo0.2 HEA has a metastable FCC single-phase microstructure. During annealing, σ phase enriched with Mo and Cr precipitates in the grains and at the grain boundaries. As the temperature increases from 700 ◦C to 1000 ◦C, the size of the precipitates grows apparently. The volume fraction of the precipitates tends to increase gradually as the annealing temperature up to 900 ◦C and then decrease at 1000 ◦C. At 800 ◦C, the volume fraction of the precipitates gradually increases as the annealing time is prolonged, eventually reaching a saturated stable value about 14%. 2. The comprehensive mechanical property of the annealed CoCrFeNiMo0.2 HEAs has a significant correspondence with the precipitates. The larger the volume fraction and the size of the precipitates, the higher the strength and the lower the plasticity of the HEA. Nearly all the annealed HEAs exhibit good strength–ductility combinations due to the significant precipitation enhancement and nanotwinning.

Author Contributions: B.L., Y.L., Q.F. and W.G. conceived and designed the experiments; C.Z. prepared the CoCrFeNiMo0.2 HEA samples and performed the microstructural characterization under the supervision of Y.L. and B.L.; H.Y. conducted the mechanical testing under the supervision of B.L. and W.G. All authors discussed the results and approved the final manuscript. Funding: This research was funded by National Natural Science Foundation of China, grant number 51771232, National Key Research and Development Plan of China, grant number 2016YFB0700302, Hunan Natural Science Foundation of China, grant number 2018JJ3477 and Special funds for future industrial development of Shenzhen, China, grant number HKHTZD20140702020004. Conflicts of Interest: The authors declare no conflicts of interest.

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