boride metal matrix composite

Accepted Manuscript Preparation and properties of medium entropy CoCrNi/boride metal matrix composite Igor Moravcik, Larissa Gouvea, Jan Cupera, Ivo Dlouhy PII:

S0925-8388(18)31074-0

DOI:

10.1016/j.jallcom.2018.03.204

Reference:

JALCOM 45434

To appear in:

Journal of Alloys and Compounds

Received Date: 20 November 2017 Revised Date:

14 March 2018

Accepted Date: 15 March 2018

Please cite this article as: I. Moravcik, L. Gouvea, J. Cupera, I. Dlouhy, Preparation and properties of medium entropy CoCrNi/boride metal matrix composite, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.03.204. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

ACCEPTED MANUSCRIPT

Preparation and properties of medium entropy CoCrNi/boride metal matrix composite Igor Moravcik, Larissa Gouvea, Jan Cupera, Ivo Dlouhy

RI PT

Institute of Materials Science and Engineering, NETME Centre, Brno University of Technology, Technicka 2896/2 Brno, Czech Republic Corresponding author: Igor Moravcik, Tel: +420 541 143 100; E-mail: [email protected]

M AN U

SC

Abstract The present work deals with synthesis, microstructural and mechanical properties’ evaluation of a metal matrix composite, produced by a combination of mechanical alloying and spark plasma sintering utilizing FCC CoCrNi medium entropy alloy as the matrix phase, with dispersion of Cr5B3 particles. Cr5B3 was formed in-situ, by the reaction of B4C, which was added to the powders prior to mechanical alloying, during sintering with Cr element present in the powders. The composite exhibited excellent tensile strength level with Rm of 1432 MPa, tensile elongation to fracture of 1.83 %, and elastic modulus of 226 GPa. Very fine dimple-like fracture morphology has been observed, despite the composite`s high strength properties.

1

TE D

Keywords: powder metallurgy, mechanical alloying, entropy, mechanical properties, metal matrix composites

Introduction

AC C

EP

The class of metallic materials, possessing near equi-atomic compositions, designated as high entropy alloys (HEA), has recently drawn significant attention of materials scientific community [1, 2]. These alloys became of great interest as a consequence of the novelty of their basic design concept. They are composed of at least five elements with very close atomic proportions, differentiating them from the traditional alloys, which usually comprise of single base element [3]. Apart only from fuelling the pure scientific curiosity to explore previously unanalysed compositions, HEAs demonstrated a significant application’s potential as they exhibited a number of extraordinary features. Some of the most intriguing properties are (but are not limited to): high fracture resistance at cryogenic temperatures [4], high tensile strength and ductility combinations [5-8], good thermal stability [9, 10], high temperature strength [11, 12] etc. A similar design approach of equi-atomic elemental proportions has been successfully applied to the alloys containing only three to four elements that were named correspondingly as medium entropy alloys (MEAs) [13-15]. Among the vast group of HEAs and MEAs bearing FCC microstructures, the single particular alloy with CoCrNi composition is standing out, owing to the remarkable combination of fracture 1

ACCEPTED MANUSCRIPT

Materials and methods

TE D

2

M AN U

SC

RI PT

resistance and strength, both being positively increased with negative increase in temperature [16]. The properties of this alloy surpassed that of extensively studied CoCrFeMnNi alloy [16]. In the last few years, the mechanistic aspects underlying the extraordinary performance of CoCrNi alloy have been extensively studied [17-20]. It has been shown that the mechanical properties are mostly governed by the formation of stress induced nano-twins (twinning induced plasticity effect - TWIP) which increase the ability of the material to accommodate plastic strain, and thus delays the onset of plastic instability and postpones catastrophic failure [7, 21]. Recently, it has been demonstrated that deformation induced, displacive transformation of FCC to HCP phase in CoCrNi alloy may take place as well, in addition to twinning [22]. In our previous research, it has been demonstrated that the CoCrNi medium entropy alloy is suitable as a matrix phase for metal matrix composites (MMC) [23], as well as its production feasibility by powder metallurgy processes. The powder metallurgy processes possess certain advantages over the casting route, seen in the simplicity of composite materials manufacturing and reduced grain sizes [23-25]. The idea of preparation of MMC based on CoCrNi composition is associated to the attempt to improve high temperature strength properties, as well as wear resistance, of this alloy by the introduction of stable ceramic reinforcement. In the light of previously mentioned evaluation, the presented work’s objective is the preparation of MMC with CoCrNi matrix and in-situ formed boride reinforcement, with the prospective use of mechanical alloying (MA) and Spark Plasma Sintering (SPS) densification processes.

AC C

EP

Elemental powders of Co, Cr and Ni (composition of CoCrNi expressed in atomic ratio) with 7.5 volume % of B4C were introduced into a, in-house made, tool steel milling bowl together with mixture of 15 mm and 20 mm bearing steel milling balls with 10:1 total ball to powder weight ratio. Elemental powders possessed industrial purity over 99 weight %. The most noticeable impurity is recognized to be Al element, present in Cr elemental powder due to manufacturing procedure. The use of powders with higher purity was intentionally avoided to simulate more accurately the conditions used in potential commercial applications. The 7.5 volume % of ceramic reinforcement was chosen to obtain reasonable trade-off between increased strength properties and reasonable ductility, based on the previous authors experience with MMC. The properties of B4C, very low density of 2.52 g cm-3, high hardness of 3800 HV and elastic modulus of 441 GPa predetermine it for the use as a reinforcement phase in MMCs, especially considering its low CTE of 4.5 ×10-6 K-1 [26]. Upon introduction into metallic matrix of CoCrNi (CTE of 17.4 ×10-6 K-1) it should produce significant strengthening upon cooling from processing temperature due to stresses arising from high CTE mismatch. The milling was performed by Pulverisette 6 planetary ball mill (Fritsch), into which the sealed milling bowl filled with a high purity nitrogen atmosphere (6.0 – 99.9999%) for powder protection was inserted; speed of 250 revolutions per minute and milling time of 35 hours was applied. To remove the powder stuck to the milling balls surfaces, toluene was further added and such wet milling has been 2

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

performed for additional 15 minutes. The mechanically alloyed powders were, then, removed from the milling bowl and dried in air. The consecutive densification has been performed in Sumitomo Coal Mining, Dr. Sinter SPS machine in Central European Institute of Technology (CEITEC). Graphite die with 30 mm inner diameter has been used. For an easy removal, graphite foils were placed between the powders and die walls. The temperature of 950 ºC with 60 MPa pressure and 10 minutes soaking time were used for the densification. Heating rate of 50 ºC min-1 has been utilized. The resulting bulk specimen was approximately 6 mm high cylinder with 30 mm diameter. Preparation of powder and bulk samples for microstructural observations was carried out by hot mounting in a polymeric resin. The samples were then mechanically ground with SiC emery papers of different particles sizes down to 1200 grit size. Subsequently, the samples were polished using diamond paste with 3 µm and 1 µm particle size. The last step of sample preparation was performed by mechano-chemical polishing, using Struers OPS suspension, to obtain perfectly flat, scratch free surfaces. X-ray diffraction analysis (XRD) of the materials phase composition was carried out using diffractometer Philips X´Pert Pro operated at 40 kV with Cu Kα. XRD patterns analysis was performed in X´Pert High Score Plus software. The scanning electron microscope (SEM) characterization of whole set of materials was performed with Carl Zeiss Ultra Plus microscope in secondary (SE) and backscattered electrons (BSE) mode respectively. Electron backscattered diffraction (EBSD) mode and energy dispersive X-ray analysis (EDS) was used for the evaluation of grain lattices and chemical composition, respectively. For EBSD measurements, the reference direction for inverse pole figures (IPF) is perpendicular to the pressing direction during SPS densification (as the texture was most likely to be formed in this direction). The same direction corresponds to X0 axis in pole figures of Fig. 5b, with X0 and Y0 axis pointing perpendicular and parallel with respect to SPS pressing direction, respectively. Both axis perpendicular to pressing direction (X0 and Z0) of Cartesian standard frame are equivalent, due to the spherical symmetry of the produced cylindrical sample. The grain boundaries have been denoted for grain misorientation angles higher than 10°. The presented values of mean grain sizes and volume fractions of respective phases have been taken as an average of measurement on three random places. Image J software has been used for image analysis of residual porosity. Measurement of Vickers hardness was carried out according to ISO 6507-1:2005, with LECO LM 274 AT hardness tester, with 0.3 kg load. The selected load was observed to be appropriate to measure the average mechanical properties of bulk with present phases within the fine grained microstructure (average diagonal of the imprint 35 µm for determined hardness level). The presented values are an average of at least 7 measurements. The tensile strength test was performed using cylindrical 28 mm long sample having a gauge length of 12.5 mm, 3.5 mm in diameter, tested with cross-head speed of 0.2 mm min-1 (i.e. strain rate of 0.25×10-4 according to EN ISO 6892) on Instron 8801 testing machine at room temperature. The elongation was measured directly on sample gauge length with clip-on extensometer. The sample was manufactured by fine turning from rectangular block with dimensions of 6 × 6 × 28 mm cut from 3

ACCEPTED MANUSCRIPT

SPSed cylinder by electric discharge machining (EDM). The value of elastic modulus has been calculated from the slope of elastic part of tensile curve.

3.1

Results

RI PT

3

CoCrNi / B4C composite powders

AC C

EP

TE D

M AN U

SC

Powder particles morphology and microstructures used for bulk material preparation, in the state after MA process, are presented in Fig. 1. Despite the use of elemental powders of Co, Cr and Ni, the MA powder particles exhibit good chemical homogeneity, as there is no visible black and white contrast in the BSE SEM mode. In addition to this, there is no evidence of layered microstructure of present element in observed as well, due to perfect elements mixing during MA processing. The dispersion of B4C has been detected by the EDS measurement in the milled powder cross section as well as small volume fraction of Cr oxides, denoted in Fig 1c, d by yellow and blue arrows. As compared to the original B4C powder (bearing of average particle size of 2 µm), the particle size after mechanical alloying has been severely reduced by the constant crushing of powder particles by milling balls during MA. This is also visible from Fig. 2 b. The XRD patterns presented in Fig. 2 a, b corresponds to blended and milled CoCrNi / B4C composite powder. While the pure elements and B4C are clearly distinguishable in the blended powder, after the MA process, only broad peaks of FCC solid solution are visible. The same phenomenon of FCC phase formation by element dissolution has been observed previously in the literature, with the peak broadening phenomena pertaining to decrease in size of crystallites by plastic deformation during MA process [8, 27]. There may be other peaks present in the milled powder pattern, however, due to extreme deformation induced to the powders, the corresponding phases are not recognizable as the background noise is significant. The peak-like formation at approximately 33° of 2 theta angle may correspond to the peak of phase with extremely reduced crystallite size, but at the same time may be just a part of the background noise.

4

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

CoCrNi / boride composite bulk

AC C

3.2

EP

Fig. 1 SEM micrographs CoCrNi / B4C mechanically alloyed powders a) powder morphology; b) powder cross-section; c), d) powder cross-section details with denoted B4C and Cr-oxide particle by yellow and blue arrows, respectively.

The XRD pattern analysis of the CoCrNi / B4C alloy mechanically alloyed powder and bulk SPSed material produced from the same powder are provided in Fig. 2 c. After the SPS densification, the major FCC solid solution phase with lattice parameter of 3.57 Å was formed. The peaks of the FCC phase are slightly shifted towards higher angles and much sharper than peaks of FCC in the milled powder. This fact contemplates to slight change in FCC phase lattice parameter and increased average crystallite size respectively, after SPS densification as compared to FCC phase detected in milled powders. Surprisingly, no traces of added B4C phase, that was observed in the milled powders prior to SPS densification, were detected in the pattern of bulk materials (Fig. 2c). Instead, the peaks of second most dominant phase were attributed to 5

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

tetragonal Cr5B3 boride phase with lattice parameter a = 5.39 Å and c = 10.17 Å. It seems that the thermodynamic stability of B4C phase in CoCrNi matrix was not sufficient, as upon high temperature sintering, boron reacted with Chromium to form previously mentioned boride. As B4C transformed to Cr5B3, these materials will be further referred to as CoCrNi / boride composite. The volume fractions of FCC and Cr5B3 boride phase calculated from diffraction pattern are 83 % and 12.4 %, respectively. There are additional peaks in the pattern conforming to the rest 4 % fraction of the microstructure, that correspond mostly to chromium carbides of Cr7C3 and Cr3C2 stoichiometry and chromium oxide phases. However, as they were impossible to precisely characterize due to their complexity and relatively small fraction on the interstice of detectability by XRD, they will be referred to as carbide and oxide phases.

Fig. 2 XRD pattern of a) CoCrNi / B4C blended powder; b) CoCrNi / B4C mechanically alloyed powder; c) sintered bulk composite. The chemical composition of the alloy is presented in Fig.3. The ratio of the elements of Co, Cr and Ni is very close to 1:1:1, as desired. It can be seen on the obtained EDS pattern that the bulk material also contains significant amounts of C and B. However, due to the very problematic determination of true C and B concentration given by the nature of the EDS method, the presented results should be taken only as a proof of B and C presence. The presence of hard 6

ACCEPTED MANUSCRIPT

EP

TE D

M AN U

SC

RI PT

ceramic particles of B4C most probably induced wear of steel milling media and milling bowl – hence the reason why elevated concentration of Fe was detected, even though it was not intentionally added to the milling bowl, prior to start of MA. The Al peaks and its concentration present in the EDS spectra originate from a trace amount of Al present in used Cr elemental powder. However, due to its significantly low concentration, no significant phase changes should be imposed to the microstructure.

AC C

Fig. 3 Chemical composition of CoCrNi / boride bulk alloy measured by SEM EDS. Due to the nature of the method, the amounts of B and C are only for reference. Fig. 4 depicts the representative microstructures of the CoCrNi / boride composite. The microstructure is not completely porous free. However, the size of the pores is rather small and their volume very low; less than 1 volume % was measured by image analysis. In addition to that, pore shape is rounded, therefore the pores should not present a sites of high stress concentration, i.e. they should not impede the material`s ductility decidedly. The microstructure is composed of FCC phase grains, with dispersed Cr5B3 boride. Additionally, the black dots (denoted by green arrow) located on the grain boundaries of all phases were determined to be carbides and oxides. The results correlate perfectly with previously presented XRD patterns results. Oxides containing mostly Cr element (as presented by EDS measurement of such particle in Fig. 5d) have been 7

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

formed as a consequence of powder contamination by oxygen – a phenomenon common to PM production methods. The carbides appear in the microstructure due to the reaction of remaining C from B4C with chromium, i.e. the free C atoms previously contained within B4C also reacted with Cr, forming a Cr carbide phase, as a consequence of Cr is the strongest carbide former among the present elements. However, as the amount of C in B4C is 4 times lower compared to B, the amount of formed carbides is relatively smaller, especially considering that the available carbon for carbide forming reaction may have been also reduced due to dissolution of carbon in FCC and Cr5B3 phases. The presence of additional nitrides, that may have formed as a consequence of the utilization of nitrogen as a protective atmosphere has not been detected i.e. the atmosphere was inert. The grain size of the major phases present in the microstructure is extremely small, as seen from Fig. 4 c. For quantitative evaluation of the microstructure, EBSD measurement has been performed with results concerning the phase distribution, orientation and grain sizes, which are presented together in Fig. 5. Average grain size of FCC phase and Cr5B3 phase is 0.77 and 0.46 µm in diameter respectively, as measured by EBSD. It is undoubtedly a consequence of the utilized rapid densification by SPS of the severely plastically deformed mechanically alloyed powders. The grain growth of FCC phase was additionally significantly impeded by present dispersed borides as well as carbides and oxides, by pinning of the traveling grain boundaries. The grains of both phases do not exhibit any type of preferential crystallographic orientation (as seen from Fig. 5a and b) in spite of the uniaxial nature of applied pressure during SPS. The grains of the Cr5B3 are located on the boundaries of the FCC grains and compose approximately 11 % of the microstructure (area fraction of Cr5B3 as measured by EBSD). Even though only 7.5 volume % of B4C was added, the reaction with Cr caused the total volume of formed Cr5B3 to be higher than 7.5 %. Due to the resolution constraints of SEM EBSD measurement, the carbides and oxide phases are not present in Fig. 5. The comprehensive analysis of boride phase by SEM, presented in Fig. 6, reveals the boride phase perfectly matches the Cr5B3 phase that was already observed in XRD patterns. The Kikuchi diffraction patterns measured from the particle in Fig. 6b closely matches the pattern of tetragonal Cr5B3 from the database, while the chemical compositions measured by point EDS and the calculated one in Fig. 6c, are almost identical. It seems some amount of Co and Ni elements is also dissolved in Cr5B3 phase, or the volume of the particle was too small for the resolution of EDS method, which caused the signal pick up from surrounding FCC grains.

8

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

Fig. 4 a), b), c) SEM micrographs in BSE mode of CoCrNi/boride bulk composite showing FCC grains denoted by yellow arrow, Cr5B3 boride particles denoted by red arrow and oxides by green arrow respectively d) EDS pattern of the oxide particle denoted by green arrow revealing high O and Cr concentration.

9

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 5 Representative SEM EBSD measurement results a) inverse pole figure orientation map, with X axis as reference direction, perpendicular to SPS pressing direction FCC grain boundaries denoted by black lines; b) pole figures of FCC and Cr5B3 phases; c) phase map showing distribution of Cr5B3 boride particles in FCC matrix; d) grain size distribution of respective phases with average grain size and volume fraction.

10

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

Fig. 6 analysis of Cr5B3 particle a) SEM micrographs with particle denoted by yellow arrow; b) corresponding electron diffraction pattern matching tetragonal Cr5B3 reflections; c) EDS point microanalysis (B content only for reference).

11

ACCEPTED MANUSCRIPT

3.3

Mechanical properties of CoCrNi / boride composite

AC C

EP

TE D

M AN U

SC

RI PT

The hardness value obtained by Vickers test was found to be 467 ± 21 HV0.3. The ultimate tensile strength (Rm) of the composite was measured at room temperature on one sample, due to the relatively limited volume of the produced material. The evaluation of results is presented in Table 1 and Fig. 7, respectively, and are also compared to previous results [18,23] of CoCrNi alloy produced by same route (combination of MA and SPS) with average grain size of 4 µm and cast alloy after high pressure torsion (HPT) and annealing with average grain size of 0.29 µm. The presence of refined grain size and ceramic particles dispersion resulted in the dramatic increase of tensile strength properties, as compared to pure powder metallurgy CoCrNi alloy from [23], as the values of both yield strength Rp0.2 and ultimate tensile strength Rm are surpassing 1400 MPa, while the ductility is severely reduced. However, very low strain hardening is also observed with the difference in the measured values of Rp0.2 and Rm reaches 7 MPa only. The values of the tensile stress are decreasing after they reach the Rm value – which may look like the presented Rm is only distinguishable yield point. However, this yield point effect occurred probably due the onset of cracking instantly after Rm as consequence of crack initiation and its rapid propagation. The total elongation value reached 1.86 % as a consequence of the pronounced strengthening by the present boride phase, as well as carbides and oxides.

12

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 7 Stress – strain curves from tensile test of CoCrNi / boride composite and pure CoCrNi alloys at room temperature.

AC C

EP

Table 1 Tensile properties of CoCrNi / boride composite in comparison with pure CrCrNi alloys. Elastic Rp0.2 Rm Elongation to Reduction in modulus E Ref. (MPa) (MPa) fracture At (%) area Z (%) (GPa) 1425

1432

226

1.86

3.9

This study

993

1100

-

20

-

[18]

1024

222

26

22

[23]

652

13

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

The results of fractographic analysis of the broken tensile test specimen are presented in Fig. 8. Fracture was initiated on the relatively large, brittle particle (as compared to the rest of microstructural features). The EDS analysis revealed that the particle is composed almost solely of pure Cr element. Most probably, some small amount of Cr particles have not been milled, thus remaining in the microstructure, despite the fact that powders were controlled for sufficient homogeneity. Such unfractured particle was not observed previously in the milled powders microstructure prior to SPS densification, as the area controlled by SEM was relatively small compared to the bulk sample volume. It is presumable that the absence of this inclusion could slightly increase the ductility of the composite. To prevent appearance of such inclusions, the milling time of the MA process will be increased in future investigation. Apart from the crack initiating inclusion, the morphology of the fracture surface is a mixture of ductile and brittle fracture behaviour. The surface is composed of very fine ductile dimples, formed by adjacent microvoids’ coalescence, but without visible inclusions and/or other particles present at the dimple bottoms. However, the fracture surface contains ridges retracting towards the fracture initiation site, i.e. morphology typical for brittle fracture behaviour. In general, considering the low value of ductility at tensile loading, the fracture should be referred to as generally brittle, with transgranular nature, often called low energy tearing.

14

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

Fig. 8 SEM micrographs of fracture surfaces of CoCrNi / boride composite tensile specimen a) fracture surface overview with crack initiation site denoted by yellow arrow; b) image of initiation site; c) image of area denoted by yellow rectangle of initiation particle interface; d) typical fracture surface morphology with fine ductile dimples.

15

ACCEPTED MANUSCRIPT

4

Discussion

AC C

EP

TE D

M AN U

SC

RI PT

The use of MA process, coupled with rapid densification by SPS, yielded powder and bulk materials respectively, with severely reduced grain sizes. The grain size reduction is attributed to the plastic deformation of the powder during the MA process, imposed by the repeated impacts of the steel milling balls [28]. In this manner, the mechanical milling presents one of the most promising severe plastic deformation (SPD) technologies for the production of advanced materials with reduced grain size, especially regarded for its simplicity [29]. It should be noted that, the B4C particle size was also decreased by powder crushing during high energy milling process – during preparation of the CoCrNi/B4C composite powder. Severely reduced crystallite size of the powders resulted in the formation of extremely grain refined microstructure after SPS densification, which were enabled by the relatively low selected sintering temperature of 950 ºC, as compared to traditional sintering [35][23][30]. The lower processing temperature and introduction of boride and oxide particles resulted in decrease of FCC phase grain size, as compared to pure CoCrNi alloy produced by the same route in last research report of the authors [23]. However, it is worth noting that the direct comparison of sintering temperature for SPS process between different studies may be sometimes slightly misleading. The methodology of the temperature measurement may be changing with different SPS machine manufacturer, and it is largely dependent also on factor such as die dimension and geometry, the type of used thermocouple, etc. The B4C particles introduced to the milling bowl with elemental powders have remained stable during the MA process. However, the presence of B4C particles in the CoCrNi alloy resulted in an in-situ reaction during sintering process, forming Cr5B3 boride phase due to increased diffusion kinetics at high temperature exposure. This phenomenon seems interesting, since one would expect B4C ceramic phase to be very stable in the metallic matrixes. The result can be explained by the calculation of enthalpy of formation change ∆Hf (calculated by Miedema`s thermodynamic model from [31, 32]). The ∆Hf for B4C is 46.5 kJ mol-1 while the value for Cr5B3 is 186.3 kJ mol-1. Since in general, more negative value of ∆Hf corresponds to higher phase stability by decreasing free energy of the system [33], the formation of Cr5B3 is evidently more thermodynamically favourable. The reaction in the solid state was additionally supported by the preceding severe mechanical activation of the powder blends during mechanical alloying process, which severely suppresses the activation energy needed for the start of boride forming reaction. During sintering process, the original B4C completely reacted as its traces have not been found in bulk materials. Considering the later preposition, it is possible to calculate the upper volume fraction of formed Cr5B3. The calculation was performed with consideration of retaining the same amounts of the present elements, with B being completely bonded to Cr, and tabulated density of Cr5B3 - 9.93 g.cm3.The rest of the microstructure was assumed to be base FCC phase partially depleted by Cr (as part of Cr atoms was partially bonded to B) and enriched by C, with density approximated by the rule of mixture to be 8.41 g.cm3. The calculated fraction is 14.6%, which correlates relatively well with the determined fraction of 11 volume % by EBSD and 12.4 % by 16

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

XRD. The slight discrepancy is most probably associated with the fact that part of present boron may be dissolved in the FCC phase (as B and C atoms possess relatively small atomic sizes, some extent of solubility in solid solution phase is expected). The remaining C atoms of original B4C additionally formed chromium carbide phase. The oxide phases present in the milled powders due to oxygen contamination remained in the bulk material accordingly, pertaining to their high characteristically thermal stability. The presence of hard particles of in-situ formed Cr5B3 boride as well as carbides and oxides caused significant increase in the ultimate tensile strength and hardness of PM CoCrNi / boride composite, as compared to previous studies [16, 23, 34]. However, the increase in strength has been accompanied by notably decreased total tensile elongation to 1.86 %. Nevertheless, the mechanical performance of presented composite is still imposing, especially compared to comparable materials produced by combination of MA and SPS, which, on occasion, does not even possess sufficient ductility to perform credible tensile test [27, 35, 36]. It may be theorized that a higher level of strength of composite microstructure should be maintained also in the elevated and high temperature range, due to dislocation and grain boundary pinning by precipitated stable ceramic particles, i.e. similar as in boride reinforced, titanium matrix composites, exhibiting increased creep resistance [12, 37, 38]. This assumption is seen as the main advantage of the CoCrNi/Boride composite as compared to highly ductile pure CoCrNi single phase alloys, and it became a central part of our future investigation on this system. The high strength of presented composite is generated by the concurrent manifestation of several strengthening phenomena. The extremely reduced average grain size of the composite, inherited from the severely plastically deformed MA powders, have caused pronounced effect of grain boundary strengthening [39]. The presence of hard borides and carbides resulted directly in the increase of strength by dislocation motion suppression (as they represent formidable obstacles for dislocation glide), by introduction of residual stresses due to CTE and elastic moduli (EM) mismatch and indirectly by grain boundary pinning effect, i.e. Zener strengthening induced by grain growth decrease during sintering [10]. As the interface between boride particles and FCC phase seems to be strong (as explained later in discussion), the contribution of load-transfer strengthening should be considered as well [41, 42]. The analysis of fracture surfaces of broken tensile test specimen demonstrated that the interface between the boride reinforcement and the matrix FCC phase exhibited very high bonding strength, as no direct interfacial decohesion has been observed in fracture surfaces. This explanation is supported by the absence of hard boride particles or other ceramic particles in the bottoms of the fracture dimples. The presumed void fracture mechanism in the composite is depicted in Fig. 9. The dimples formed predominantly at places possessing the highest stress triaxiality (grain boundaries, intersection of slip systems, etc.) by ductile tearing of FCC phase, i.e. not by FCC / boride interphase boundary decohesion. Presumably, the dimples had nucleated and subsequently grew between narrowly spaced angular boride particles by matrix voiding microprocess (displayed in inset of Fig. 9), same as previously observed in [43]. At such locations, the local stress raised by the angular edges of boride particles, acting as stress concentrators, prevails over the cohesive strength of the FCC phase and new free surface, in the 17

ACCEPTED MANUSCRIPT

EP

TE D

M AN U

SC

RI PT

form of void, is created. In this manner, transgranular fracture occurred throughout fine FCC grains. Therefore, it is interesting to note that despite relatively high ceramic loading, the fracture process is governed by the plastic deformation of the matrix (CoCrNi) FCC phase. The strength of the plastic FCC phase appears to be lower than cohesive strength of the interphase boundary between boride and FCC phase. Therefore, the interphase boundary between in-situ formed Cr5B3 borides must be inherently strong, which is one of the basic requirements for ductile composites with high ceramic loadings [44]. It should be noted that, in the present study, the fracture process was initiated on brittle Cr inclusion particle. To avoid such inclusions and further improve properties, a simple increase in the milling time could be applied in the future preparation.

AC C

Fig. 9 Mechanism of transgranular crack formation in CoCrNi / boride composite by adjacent microvoids coalescence on the crack tip in FCC matrix phase induced by stress concentration between rounded boride corners. The presented results raise the question concerning the effectivity of production of the medium entropy alloy / boride in-situ composite. It has been widely accepted that HEAs and MEAs are promising materials for the employment as matrix phase [44-46]. However, the formation of insitu boride phase by reaction of Cr with B4C is more disputable. It may seem that, it would be more efficient to produce boride reinforced composite by conventional, ex-situ route, i.e. by the boride reinforcement phase directly introduced to melted metal matrix and it would not transform during manufacturing process. However, in-situ composites can be more cost effective and deliver higher mechanical strength, particularly at relatively higher operating temperature [47].

18

ACCEPTED MANUSCRIPT

5

M AN U

SC

RI PT

This is usually attributed to better mechanical bonding between matrix and reinforcement. The presented results from this contribution are in good agreement with this claim. Another aspect is the use of B4C phase as a precursor phase for boride formation. It is more usual to either use precursor pure elements, or oxides, to fuel reinforcement forming reactions [46]. The utilization of pure B as a precursor for boride formation is however limited by its extremely high price, while the use of cheap boron oxide (e.g. B2O3) is associated with appearance of significant amount of undesirable oxide phases [47, 48], sometimes surpassing the total volume of boride reinforcement (given by stoichiometry of boron oxide ). B4C is relatively cheap and it has been already successfully utilized as a precursor phase for the production of TiB2/TiC HEA coatings [49] as well as bulk materials [48, 50]. Therefore, its use as a raw material is sufficiently justified. As such, the production of boride reinforcement by decomposition of B4C should be in theory achievable in any metallic matrix containing corresponding fraction of element with high boron chemical affinity.

Conclusions

The preparation of in-situ CoCrNi / boride composite by a combination of mechanical alloying and spark plasma sintering has been carried out, with subsequent microstructural and mechanical properties’ characterization. The conclusions are drawn as follows:

• • •

TE D

•

EP

•

Metal matrix composites with extremely fine grain size may be readily prepared by the combination of mechanical alloying and spark plasma sintering; B4C ceramic phase exhibited significant instability during high temperature sintering in Cr- rich metal matrix; The introduction of B4C resulted in the in-situ solid state reaction during sintering, producing 11 volume % of Cr5B3 boride phase; The CoCrNi/boride in-situ composite possessed high strength values, surpassing ultimate tensile strength Rm of 1400 MPa and tensile elongation of 1.86 %; Ductile dimples have been observed on fracture surfaces; Cr5B3 boride reinforcement and FCC CoCrNi matrix phase exhibited especially high interfacial strength.

AC C

•

Acknowledgments The research was co-funded by the Ministry of Education, Youth and Sports within the „National Sustainability Programme I” (NETME CENTRE PLUS - LO1202). Authors would also like to thank Dr. David Salomon and Hua Tan, MSc. from CEITEC for their help with the SPS sample’s densification experiments. The presented results have been obtained within the PhD. thesis project of the corresponding author [51].

19

ACCEPTED MANUSCRIPT

RI PT

References:

AC C

EP

TE D

M AN U

SC

[1] B.D. Miracle, D.J. Miller, N.O. Senkov, C. Woodward, D.M. Uchic, J. Tiley, Exploration and Development of High Entropy Alloys for Structural Applications, Entropy 16(1) (2014). [2] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Progress in Materials Science 61(0) (2014) 1-93. [3] B.S. Murty, J.W. Yeh, S. Ranganathan, High-Entropy Alloys, Elsevier, ButterworthHeinemann, London, 2014. [4] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, A fractureresistant high-entropy alloy for cryogenic applications, Science 345(6201) (2014) 1153-1158. [5] W.H. Liu, Z.P. Lu, J.Y. He, J.H. Luan, Z.J. Wang, B. Liu, Y. Liu, M.W. Chen, C.T. Liu, Ductile CoCrFeNiMox high entropy alloys strengthened by hard intermetallic phases, Acta Materialia 116 (2016) 332-342. [6] J.Y. He, H. Wang, H.L. Huang, X.D. Xu, M.W. Chen, Y. Wu, X.J. Liu, T.G. Nieh, K. An, Z.P. Lu, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Materialia 102 (2016) 187-196. [7] D. Raabe, C.C. Tasan, H. Springer, M. Bausch, From High-Entropy Alloys to High-Entropy Steels, steel research international 86(10) (2015) 1127-1138. [8] I. Moravcik, J. Cizek, J. Zapletal, Z. Kovacova, J. Vesely, P. Minarik, M. Kitzmantel, E. Neubauer, I. Dlouhy, Microstructure and mechanical properties of Ni1,5Co1,5CrFeTi0,5 high entropy alloy fabricated by mechanical alloying and spark plasma sintering, Materials & Design. [9] M.J. Yao, K.G. Pradeep, C.C. Tasan, D. Raabe, A novel, single phase, non-equiatomic FeMnNiCoCr high-entropy alloy with exceptional phase stability and tensile ductility, Scripta Materialia 72–73 (2014) 5-8. [10] O.N. Senkov, J.M. Scott, S.V. Senkova, F. Meisenkothen, D.B. Miracle, C.F. Woodward, Microstructure and elevated temperature properties of a refractory TaNbHfZrTi alloy, Journal of Materials Science 47(9) (2012) 4062-4074. [11] M.A. Manzoni, S. Singh, M.H. Daoud, R. Popp, R. Völkl, U. Glatzel, N. Wanderka, On the Path to Optimizing the Al-Co-Cr-Cu-Fe-Ni-Ti High Entropy Alloy Family for High Temperature Applications, Entropy 18(4) (2016). [12] H.M. Daoud, A.M. Manzoni, N. Wanderka, U. Glatzel, High-Temperature Tensile Strength of Al10Co25Cr8Fe15Ni36Ti6 Compositionally Complex Alloy (High-Entropy Alloy), JOM 67(10) (2015) 2271-2277. [13] K.V.S. Thurston, B. Gludovatz, A. Hohenwarter, G. Laplanche, E.P. George, R.O. Ritchie, Effect of temperature on the fatigue-crack growth behavior of the high-entropy alloy CrMnFeCoNi, Intermetallics 88 (2017) 65-72. [14] Z. Wu, H. Bei, F. Otto, G.M. Pharr, E.P. George, Recovery, recrystallization, grain growth and phase stability of a family of FCC-structured multi-component equiatomic solid solution alloys, Intermetallics 46 (2014) 131-140. [15] A. Gali, E.P. George, Tensile properties of high- and medium-entropy alloys, Intermetallics 39 (2013) 74-78. 20

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

[16] B.H. Gludovatz, A. Thurston, KV. Bei. H. Wu. Z. George, EP. Ritchie, RO., Exceptional damage-tolerance of a medium-entropy alloy CrCoNi at cryogenic temperatures, Nature communications, 2016. [17] Y.L. Zhao, T. Yang, Y. Tong, J. Wang, J.H. Luan, Z.B. Jiao, D. Chen, Y. Yang, A. Hu, C.T. Liu, J.J. Kai, Heterogeneous precipitation behavior and stacking-fault-mediated deformation in a CoCrNi-based medium-entropy alloy, Acta Materialia 138 (2017) 72-82. [18] S. Yoshida, T. Bhattacharjee, Y. Bai, N. Tsuji, Friction stress and Hall-Petch relationship in CoCrNi equi-atomic medium entropy alloy processed by severe plastic deformation and subsequent annealing, Scripta Materialia 134 (2017) 33-36. [19] Z. Zhang, H. Sheng, Z. Wang, B. Gludovatz, Z. Zhang, E.P. George, Q. Yu, S.X. Mao, R.O. Ritchie, Dislocation mechanisms and 3D twin architectures generate exceptional strengthductility-toughness combination in CrCoNi medium-entropy alloy, Nature Communications 8 (2017) 14390. [20] G. Laplanche, A. Kostka, C. Reinhart, J. Hunfeld, G. Eggeler, E.P. George, Reasons for the superior mechanical properties of medium-entropy CrCoNi compared to high-entropy CrMnFeCoNi, Acta Materialia 128 (2017) 292-303. [21] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, High manganese austenitic twinning induced plasticity steels: A review of the microstructure properties relationships, Current Opinion in Solid State and Materials Science 15(4) (2011) 141-168. [22] J. Miao, C.E. Slone, T.M. Smith, C. Niu, H. Bei, M. Ghazisaeidi, G.M. Pharr, M.J. Mills, The evolution of the deformation substructure in a Ni-Co-Cr equiatomic solid solution alloy, Acta Materialia 132 (2017) 35-48. [23] I. Moravcik, J. Cizek, Z. Kovacova, J. Nejezchlebova, M. Kitzmantel, E. Neubauer, I. Kubena, V. Hornik, I. Dlouhy, Mechanical and microstructural characterization of powder metallurgy CoCrNi medium entropy alloy, Materials Science and Engineering: A 701 (2017) 370-380. [24] G. Fan, R. Xu, Z. Tan, D. Zhang, Z. Li, Development of Flake Powder Metallurgy in Fabricating Metal Matrix Composites: A Review, Acta Metallurgica Sinica (English Letters) 27(5) (2014) 806-815. [25] C.C. Eiselt, H. Schendzielorz, A. Seubert, B. Hary, Y. de Carlan, P. Diano, B. Perrin, D. Cedat, ODS-materials for high temperature applications in advanced nuclear systems, Nuclear Materials and Energy 9 (2016) 22-28. [26] H. Holleck, Material selection for hard coatings, J. Vac. Sci. Technol. A 4, 2661 1986. [27] Z. Fu, W. Chen, H. Xiao, L. Zhou, D. Zhu, S. Yang, Fabrication and properties of nanocrystalline Co0.5FeNiCrTi0.5 high entropy alloy by MA–SPS technique, Materials & Design 44 (2013) 535-539. [28] J. Joardar, S.K. Pabi, B.S. Murty, Milling criteria for the synthesis of nanocrystalline NiAl by mechanical alloying, Journal of Alloys and Compounds 429(1–2) (2007) 204-210. [29] X. Zhang, H. Wang, M. Kassem, J. Narayan, C.C. Koch, Preparation of bulk ultrafinegrained and nanostructured Zn, Al and their alloys by in situ consolidation of powders during mechanical attrition, Scripta Materialia 46(9) (2002) 661-665. [30] I. Sulima, P. Putyra, P. Hyjek, T. Tokarski, Effect of SPS parameters on densification and properties of steel matrix composites, Advanced Powder Technology 26(4) (2015) 1152-1161. [31] A.K. Niessen, F.R. de Boer, R. Boom, P.F. de Châtel, W.C.M. Mattens, A.R. Miedema, Model predictions for the enthalpy of formation of transition metal alloys II, Calphad 7(1) (1983) 51-70. 21

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

[32] R.F. Zhang, S.H. Zhang, Z.J. He, J. Jing, S.H. Sheng, Miedema Calculator: A thermodynamic platform for predicting formation enthalpies of alloys within framework of Miedema’s Theory, Computer Physics Communications 209(Supplement C) (2016) 58-69. [33] X. Yang, Y. Zhang, Prediction of high-entropy stabilized solid-solution in multi-component alloys, Materials Chemistry and Physics 132(2–3) (2012) 233-238. [34] Z. Wu, H. Bei, G.M. Pharr, E.P. George, Temperature dependence of the mechanical properties of equiatomic solid solution alloys with face-centered cubic crystal structures, Acta Materialia 81 (2014) 428-441. [35] S. Fang, W. Chen, Z. Fu, Microstructure and mechanical properties of twinned Al0.5CrFeNiCo0.3C0.2 high entropy alloy processed by mechanical alloying and spark plasma sintering, Materials & Design (1980-2015) 54 (2014) 973-979. [36] C. Wang, W. Ji, Z. Fu, Mechanical alloying and spark plasma sintering of CoCrFeNiMnAl high-entropy alloy, Advanced Powder Technology 25(4) (2014) 1334-1338. [37] H.T. Tsang, C.G. Chao, C.Y. Ma, Effects of volume fraction of reinforcement on tensile and creep properties of in-situ TiBTi MMC, Scr. Mater. 37 (1997) 1359–1365. doi:10.1016/S1359-6462(97)00251-0. [38] K.S. Ravi Chandran, K.B. Panda, S.S. Sahay, TiBw-reinforced Ti composites: Processing, properties, application prospects, and research needs, JOM. 56 (2004) 42–48. doi:10.1007/s11837-004-0127-1. [39] N. Chawla, Y.L. Shen, Mechanical Behavior of Particle Reinforced Metal Matrix Composites, Advanced Engineering Materials 3(6) (2001) 357-370. [40] A.F. Mirza, L.D. Chen, A Unified Model for the Prediction of Yield Strength in ParticulateReinforced Metal Matrix Nanocomposites, Materials 8(8) (2015). [41] M. Kouzeli, A. Mortensen, Size dependent strengthening in particle reinforced aluminium, Acta Materialia 50(1) (2002) 39-51. [42] A. Sanaty-Zadeh, Comparison between current models for the strength of particulatereinforced metal matrix nanocomposites with emphasis on consideration of Hall–Petch effect, Materials Science and Engineering: A 531 (2012) 112-118. [43] A. Miserez, R. Müller, A. Rossoll, L. Weber, A. Mortensen, Particle reinforced metals of high ceramic content, Materials Science and Engineering: A 387–389 (2004) 822-831. [44] J.Y. He, W.H. Liu, H. Wang, Y. Wu, X.J. Liu, T.G. Nieh, Z.P. Lu, Effects of Al addition on structural evolution and tensile properties of the FeCoNiCrMn high-entropy alloy system, Acta Materialia 62 (2014) 105-113. [45] G. Zhu, Y. Liu, J. Ye, Fabrication and properties of Ti(C,N)-based cermets with multicomponent AlCoCrFeNi high-entropy alloys binder, Materials Letters 113 (2013) 80-82. [46] C.-M. Lin, C.-W. Tsai, S.-M. Huang, C.-C. Yang, J.-W. Yeh, New TiC/Co1.5CrFeNi1.5Ti0.5 Cermet with Slow TiC Coarsening During Sintering, JOM 66(10) (2014) 2050-2056. [47] L. Lü, M.O. Lai, Y. Su, H.L. Teo, C.F. Feng, In situ TiB2 reinforced Al alloy composites, Scripta Materialia 45(9) (2001) 1017-1023. [48] H. Zhao, Y.-B. Cheng, Formation of TiB2–TiC composites by reactive sintering, Ceramics International 25(4) (1999) 353-358. [49] J. Cheng, D. Liu, X. Liang, Y. Chen, Evolution of microstructure and mechanical properties of in situ synthesized TiC–TiB2/CoCrCuFeNi high entropy alloy coatings, Surface and Coatings Technology 281(Supplement C) (2015) 109-116. 22

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

[50] G.J. Zhang, Z.Z. Jin, X.M. Yue, Reaction synthesis of TiB2-SiC composites from TiH2-SiB4C, Materials Letters 25(3) (1995) 97-100. [51] I. Moravcik, METAL MATRIX COMPOSITES PREPARED BY POWDER METALLURGY ROUTE, PhD thesis, Brno University of Technology 2017, 149p. https://www.vutbr.cz/www_base/zav_prace_soubor_verejne.php?file_id=160762

23

ACCEPTED MANUSCRIPT

Research highlights Preparation and properties of medium entropy CoCrNi/boride metal matrix composite

RI PT

SC M AN U

•

TE D

•

EP

•

CoCrNi medium entropy alloy metal matrix composite has been produced for the first time Valid tensile test has been performed on material produced by mechanical alloying and spark plasma sintering the composite exhibited excellent tensile strength over 1400MPa with 1.86 % total elongation ductile fracture mode with dimple morphology has been observed on fracture surfaces

AC C

•