Columnar to equiaxed transition and grain refinement of cast CrCoNi medium-entropy alloy by microalloying with titanium and carbon

Accepted Manuscript Columnar to equiaxed transition and grain refinement of cast CrCoNi medium-entropy alloy by microalloying with titanium and carbon X.W. Liu, G. Laplanche, A. Kostka, S.G. Fries, J. Pfetzing-Micklich, G. Liu, E.P. George PII:

S0925-8388(18)33861-1

DOI:

10.1016/j.jallcom.2018.10.187

Reference:

JALCOM 48005

To appear in:

Journal of Alloys and Compounds

Received Date: 6 September 2018 Revised Date:

13 October 2018

Accepted Date: 15 October 2018

Please cite this article as: X.W. Liu, G. Laplanche, A. Kostka, S.G. Fries, J. Pfetzing-Micklich, G. Liu, E.P. George, Columnar to equiaxed transition and grain refinement of cast CrCoNi medium-entropy alloy by microalloying with titanium and carbon, Journal of Alloys and Compounds (2018), doi: https:// doi.org/10.1016/j.jallcom.2018.10.187. 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.

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Columnar to equiaxed transition and grain refinement of cast CrCoNi mediumentropy alloy by microalloying with titanium and carbon

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X.W. Liua,b,c*, G. Laplanchea, A. Kostkab, S.G. Friesd, J. Pfetzing-Micklichb, G. Liua,b, E.P. Georgee,f*

Institute for Materials, Ruhr-University Bochum, D-44801 Bochum, Germany

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Materials Research Department and Center for Interface Dominated Materials (ZGH),

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a

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Ruhr-University Bochum, D-44801 Bochum, Germany

State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, Wuhan 430074, China

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Interdisciplinary Centre for Advanced Materials Simulation, Ruhr-University Bochum,

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Bochum 44801, Germany

Oak Ridge National Laboratory, Materials Science and Technology Division, Oak

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Ridge, TN 37831, USA

University of Tennessee, Materials Science and Engineering Department, Knoxville,

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TN 37996, USA

*Corresponding authors: [email protected] (X.W. Liu), [email protected] (E.P. George)

Abstract: Thermomechanical processing has been used to control the grain size/shape of the equiatomic CrCoNi medium-entropy alloy (MEA) and obtain excellent strength and ductility. However, in the cast state, the alloy has coarse columnar grains with average widths and lengths of approximately 120 and 1000 µm, respectively, resulting in inferior

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mechanical properties. To overcome this deficiency, here we microalloyed with Ti and C and successfully changed the grain shape (from columnar to equiaxed) and refined the

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grain size. The degree to which the microstructure changes depends on the amount of Ti and C added, with the best results obtained at 0.4 at.% each. In the optimal alloy [(CrCoNi)99.2Ti0.4C0.4], the as-cast grains were nearly equiaxed with a uniform size of ~75

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µm. Associated with this change in grain shape/size was a significant improvement of yield strength, ultimate tensile strength and elongation to fracture at both 293 and 77 K.

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The columnar to equiaxed transition is attributed to the strong mutual affinity of C and Ti, which leads to their build-up ahead of the solid-liquid interface and, in turn, to enhanced constitutional undercooling.

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Keywords: Medium and high entropy alloys; Grain refinement; Cast microstructure;

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Constitutional undercooling; Mechanical properties

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1. Introduction The CrMnFeCoNi high-entropy alloy (HEA) and its medium-entropy subsystems

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with the face-centered cubic (FCC) structure (also referred to as multi-principal-element alloys in the literature) have attracted extensive attention during the past decade and a half (e.g. [1-28]). Among these alloys, the CrCoNi medium-entropy alloy (MEA) exhibits

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the best strength-ductility/toughness combination at cryogenic temperatures [7, 14, 19, 29]. It can therefore serve as a base for the development of promising engineering alloys

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in the future.

For nearly all engineering alloys, casting is an important step that enables the production of complex-shape components. Typically, cast alloys exhibit coarse and

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anisotropic solidification microstructures with grains aligned along the heat-flow direction. This produces bulky columnar structures outside and a relatively small region of coarse equiaxed grains in the center. The coarseness of the microstructure increases as

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the solidification/cooling rate decreases (i.e., as the size of the casting increases). Coarse

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columnar microstructures form also in cast HEAs [30] and are generally undesirable from a mechanical properties standpoint. Cast microstructures can be broken down by deformation processing and recrystallization after which the tensile properties of at least some HEAs are rather good, as was shown for the first time with CrFeCoNi and CrMnFeCoNi [2, 3]. Grain refinement in HEAs results in classical Hall-Petch strengthening [2, 5, 31] as well as simultaneous increases in strength and ductility [32, 33]. 4

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While wrought alloys can benefit from grain refinement during processing, cast alloys offer limited opportunities to manipulate the grain morphology and refine the grain

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size after the castings are made (especially if they have complex shapes). Consequently, the as-cast grain size, to a large extent, determines the final properties. It is important, therefore, to refine the as-cast grain size, not only to improve mechanical properties but

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also to decrease the tendency for hot tearing and to have a more dispersed and refined porosity distribution after solidification. The columnar to equiaxed transition (CET) and

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grain refinement of cast high- and medium-entropy alloys have received little attention so far [34]. Consequently, the methods and mechanisms that can be deployed in such materials to alter their as-cast grain size/shape remain unclear.

There are several methods for refining the as-cast microstructures such as

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inoculation/heterogeneous nucleation, rapid solidification, application of external fields (e.g. mechanical, electromagnetic, ultrasonic and stirring (e.g. [35-37]). Some of these are

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difficult to implement in the generally used laboratory-scale processes, which typically involve induction/arc melting and drop casting in a closed vacuum furnace. In these

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cases, alternative approaches using intrinsic grain refinement methods are needed. Heterogeneous nucleation has long been known to be a powerful method for grain refinement and has been used in foundries for years. Another method involves the use of certain solutes that segregate during solidification and has been shown to be an effective way to reduce grain sizes in various alloys [38-42]. In the latter case, the notion is that growth-restricting solutes induce constitutional undercooling, the extent of which depends on the specific species and amounts, and the undercooling promotes grain 5

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refinement. Here, we selected the CrCoNi alloy as a model MEA to investigate the CET and grain refinement via trace additions of Ti and C. Our results show that it is indeed

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possible to achieve significant grain refinement by the addition of optimum levels of Ti and C and the resulting fine-grained MEA has superior tensile properties. 2. Experimental

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Alloys with compositions (CrCoNi)100-2xTixCx in at.%, where x=0–1 (denoted as

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TixCx MEAs hereafter) were prepared by arc melting a mixture of the pure metals Cr, Co, Ni and Ti (purity >99.9 wt.%) and Cr3C2 powder (purity >99 wt.%) in a high-purity argon atmosphere. After the raw materials were loaded in the arc melter, the chamber was first evacuated to ~5 × 10-2 Pa and then backfilled with pure Ar. This process was repeated

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twice. After the final backfill, the Ar pressure in the chamber was ~7 × 104 Pa. Before the alloys were melted, a small piece of pure Ti was melted to getter any residual oxygen or nitrogen that might be present in the chamber. The buttons were flipped and remelted five

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times to ensure complete melting and through mixing of the raw materials in the liquid

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before drop-casting into a rectangular copper mold with dimensions of 8 × 17 × 65 mm3. To examine the cast microstructures, the ingot cross sections were ground with SiC

paper to a grit size of 7 µm and polished with diamond suspension down to 1 µm. A final polishing step was performed using a vibratory polisher (Buehler Vibromet 2) and colloidal silica with a particle size of 0.06 µm for 48 h. The microstructures were examined in a FEI Quanta FEG 650 field emission gun scanning electron microscope equipped with electron backscatter diffraction (EBSD) at an accelerating voltage of 30 6

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kV. The EBSD data were analyzed with the OIM Analysis software (version 6.2.0), following the procedure in [8]. We focused on the lower parts of the cast ingots, which

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tend to have more stable solidification conditions compared to regions near the entrance of the mold. Grain sizes were measured with the linear intercept method outlined in ASTM E112-10. In this work, equiaxed grains were defined as those with an aspect ratio

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less than 2, while those with larger aspect ratios were defined as columnar grains; aspect ratio was determined by taking the ratio of the longest to smallest dimension [43]. Phase

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characterization was carried out by X-ray diffraction using a PANalytical X’Pert Pro MRD diffractometer (Cu Kα radiation λ = 0.154 nm; 2θ-range from 20° to 120°; step size ∆2θ = 0.006°; integration time 280 s).

Thin sheets with a thickness of ~400 µm were extracted from the as-cast specimens

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and ground down to foils with a thickness of about ~100 µm. Discs with a diameter of 3 mm were punched out of the foils and electron transparent regions were produced in a

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twin-jet polishing system using an electrolyte consisting of 70 vol.% methanol, 20 vol.% glycerine, and 10 vol.% perchloric acid at a temperature of 258 K and an applied voltage

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of 25 V. TEM investigations were conducted in an FEI Tecnai F20 G2 S-Twin equipped with a field emission gun (FEG) and a high angle annular dark field (HAADF) detector. As described later, the Ti0.4C0.4 alloy exhibited the finest equiaxed grains in the as-

cast state. Therefore, it was chosen for measurement of tensile properties, along with the columnar grained CrCoNi MEA for comparison. Flat dog-bone shaped tensile specimens with a gauge length of 20 mm were machined from the cast ingots by electric discharge machining with their longitudinal axes parallel to the mold-filling direction. All the 7

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specimens were obtained from regions near the ingot surface and away from the central axis of the ingots where some porosity was usually present. The tensile specimens were

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ground carefully on each side through 1000-grit SiC paper, resulting in a final specimen thickness of ~1.2 mm and a gauge section width of ~4 mm. Vickers microhardness indents spaced 0.5 mm apart were made along the specimen gauge lengths using a KB30

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BVZ Vickers Hardness tester (KB Prüftechnik GmbH, Germany) with a force of 400 g to enable the calibration of strain by means of an optical traveling microscope, similar to the

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process described in an earlier paper [19]. Tensile tests were performed at an engineering strain rate of 10-3 s-1 in a screw-driven Zwick/Roell Z100 test rig at 77 K and 293 K (room temperature, RT). At least two tests using samples from different parts of the ingots were performed for each alloy at each temperature. For tests at 77 K, the specimen

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and grips were immersed in a liquid nitrogen bath for about 10–15 min before the start of the test and they stayed completely submerged during the test. Room-temperature tests

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were performed in air.

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3. Results and discussion

Fig. 1a is a schematic diagram showing a cross-section located 30 mm from the

bottom of the casting where microstructural examination was performed. On this section, microstructures were examined both close to the mold-ingot interface [red square labeled (b)] and near the center of the cast ingot [blue square labeled (c)]. EBSD images of ascast TixCx MEAs taken at these locations are shown in Fig. 1 b and c, respectively. Close to the mold (Fig. 1b), the base CrCoNi MEA exhibits anisotropic columnar grains with 8

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average lengths of ~1000 µm and widths of ~120 µm nearly perpendicular to the surface of the copper mold along the heat transfer direction. Upon addition of Ti and C, a shape

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transition from columnar to equiaxed occurs with the number of equiaxed grains increasing at first and then decreasing with increasing Ti/C additions. The Ti0.4C0.4 MEA shows the finest, uniformly distributed, equiaxed grains with a mean grain size of about

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75 µm. As mentioned before, tensile specimens were machined from these regions close to the mold interface (to avoid the last regions to solidify near the center where there was

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some porosity). The relevant grain sizes of these specimens were considered to be ~120 µm (the widths of the columnar grains in the base alloy) and ~75 µm (size of the equiaxed grains in the Ti0.4C0.4 alloy).

For comparison, the grain size in the center of the cast bars (Fig. 1c) averaged over

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an area of around 3 × 3 mm2 was measured, and the results are shown in Fig. 2. Again, the finest grain size is obtained for Ti and C additions of 0.4 at.% each. The

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microstructures of the MEAs at a distance of 10 mm from the bottom, which is expected to have a faster cooling rate, were also examined (Fig. 3) and found to be approximately

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similar to those in Fig. 1 (to save space, only Ti0.2C0.2, Ti0.4C0.4 and Ti0.5C0.5 MEAs are shown in Fig. 3). The main differences are that the Ti0.2C0.2 MEA shows nearly fully equiaxed grains with the finest grain size of ~65 µm, the CrCoNi MEA has even coarser grains, and the Ti0.5C0.5 MEA exhibits a mixture of columnar and some equiaxed grains. However, the Ti0.4C0.4 MEA exhibits nearly the same microstructures as in Fig. 1. Based on these results, it appears that the optimum Ti/C concentrations for grain refinement shift to lower levels with increasing cooling rate. In the remainder of this paper we focus 9

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on the Ti0.4C0.4 MEA, denoted as EG (equiaxed grains) MEA, and compare it with the base CrCoNi, denoted as CG (columnar grains) MEA.

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Fig. 4 shows representative back scattered electron images of the CG and EG MEAs taken in the SEM along with corresponding elemental EDX maps. The X-ray maps show dendritic morphology inside the grains of both the CG and EG MEAs (easier to discern in

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the larger grains of the latter). The dendrites are enriched in Co, the inter-dendritic regions are enriched in Cr, and Ni seems relatively uniformly distributed (especially in

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the CG MEA). These results are different from what has been previously reported for the CrMnFeCoNi HEA [10], namely that the elements with the highest melting points are segregated in the dendritic cores. In the EG MEA, in addition to Cr, Ti is also enriched in the interdendritic regions. No obvious segregation of C was observed, because of the low

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EDX sensitivity for light elements. A compositional scan was conducted along the black line in Fig. 4b and the results for Cr, Co, Ni and Ti are shown in Fig. 5. For clarity, the

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locations where the black line intersects interdendritic regions are marked with red lines in Fig. 5. The concentration profiles in Fig. 5 clearly show that the interdendritic regions

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are slightly enriched in Ti (up to 0.7 at.%) and Cr and depleted in Co while the Ni concentration stays roughly constant through the grains and across the grain boundaries. Such differences between dendritic and interdendritic regions are often seen in castings, including in high-entropy alloys (e.g., [10]), and originate from the different compositions of the liquid and solid at different temperatures. To eliminate this casting segregation (“coring”) in high- and medium-entropy alloys, long-term homogenization at high temperatures is usually carried out in the solid state (e.g., [2, 3, 7]). 10

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The pseudo-binary phase diagram of (CrCoNi) – Ti/C system was calculated by ThermoCalc with TCFE7 database and the results are shown in Fig. 6a. For the Ti0.4C0.4 MEA, the equilibrium phases are FCC + TiC (labeled as Fcc_A1 and Fcc_A1#2) above

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about 950 oC and FCC + TiC + M23C6 at lower temperatures to about 875 oC. However, XRD analysis revealed only FCC peaks (Fig. 6b). Since phases with a low volume

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fraction cannot be detected by XRD, further analysis by TEM was conducted. Representative grain boundaries (GBs) near a triple point in the EG MEA are shown in

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Fig. 6c (imaged in the BSE mode). In addition to smooth GBs (white arrows), serrated GBs were commonly seen (red arrow). This is in sharp contrast to the CG alloy where only smooth GBs were visible (not shown here). A TEM bright-field (BF) image of a serrated GB is shown in Fig. 6d, where the inset shows the selected area diffraction

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(SAD) pattern from the circled serration. SAD analysis indexed it to be M23C6. No TiC was found either in the intergranular or intragranular regions by TEM analysis. Only a

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few isolated spherical TiO2 particles with a size of several hundred nanometers were detected. The calculated phase diagram in Fig. 6a predicts that further secondary phases

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(e.g. two intermetallic phases: Ni3Ti with hexagonal close-packed superlattice [44] and space group PM-3M [45] and a sigma phase with topologically close-packed structure and space group P42  [46], and two solid solutions: a hexagonal phase and a BCC phase) should be stable at low temperatures. These secondary phases were not detected in the present study which may indicate that our as-cast MEA is in a metastable state. In the related CrMnFeCoNi and CrFeCoNi alloys, second-phase precipitates have been observed after prolonged anneals at intermediate/low temperatures [13, 17, 47]. Further 11

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work is needed to investigate phase stability in this alloy and to determine whether the phases predicted in the calculated phase diagram (Fig. 6a) will in fact form after

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annealing. Fig. 7a shows the engineering stress–engineering plastic strain curves of the CG and EG MEA at 293 and 77 K. At both test temperatures, the yield strength, ultimate

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tensile strength and elongation to fracture of the (CrCoNi)99.2Ti0.4C0.4 alloy are superior to those of the base CrCoNi alloy. Furthermore, similar to the base CrMnFeCoNi and

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CrCoNi alloys containing no added Ti and C [2, 6, 7, 14, 19], the tensile properties at 77 K are superior to those at 293 K, likely due to easier nanoscale twinning at lower temperatures. Consistent with this, the strain hardening rates for both MEAs at 77 K are

(Fig. 7c).

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higher than at 293 K and those of the EG MEA are higher than those of the CG MEA

The results of this work show clearly that microalloying with Ti and C can induce

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a columnar to equiaxed transition and grain refinement in the CrCoNi MEA under the current solidification conditions with a concomitant improvement of its tensile properties.

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One striking feature is the simultaneous enhancement of strength and elongation at both tested temperatures, which is difficult to obtain by other hardening methods, e.g. precipitation and dislocation hardening [21, 48]. Grain refinement in the tens to hundreds of micrometers range is one of the few ways in which strength can be increased along with ductility and we believe that is the operative mechanism here. In the Hall-Petch relationship, σ = + ⁄ , the lattice friction ( ) and strengthening coefficient for the CrCoNi MEA have been previously determined to be 216 MPa and 568 MPa µm1/2, 12

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respectively [21]. Considering grain sizes of ~120 µm for the CG MEA (width of the columnar grains) and ~75 µm for the EG MEA, their RT yield stresses should be ~268

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MPa and ~282 MPa, respectively. However, the values measured in this study are ~230 MPa and ~302 MPa, indicating that texture may play a role as well, which needs to be systematically investigated in the future. Moreover, it is not clear at present how Ti and C

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additions affect the lattice friction. The improved work hardening is ascribed to the grain refinement and associated dislocation–grain boundary interactions [49, 50].

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The columnar to equiaxed transition and grain refinement induced by Ti and C may result from heterogeneous nucleation at particles (given that the solidification conditions were otherwise controlled to be very similar for each sample). TiC is one such possible nucleant since it is stable at high temperatures. However, TiC particles were not detected

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in our TEM analyses, either on grain boundaries or in grain interiors, possibly because their levels were below their individual solubilities [51, 52] consistent with previous

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observations in many alloys [51-53]. Wu et al. [54] added 0.5 at.% C (similar to the concentration in our alloy) to the CrMnFeCoNi HEA and obtained a single phase after

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recrystallization at 1100 oC for 1 h. Stepanov et al. [55] found an insignificant fraction of nanoscale carbides in the (CoCrFeNiMn)98C2 HEA (which has about five times more C than our alloy) in the as-cast state. Although the exact solubility of Ti and C in CrCoNi is not known, from the calculated pseudo-binary CrCoNi-Ti and CrCoNi-C phase diagrams (Fig. 8), both elements have substantial solubility (>0.5 at.%) in FCC matrices at high temperature; during the relatively fast cooling of our small castings in copper molds, it may be possible to trap them in solid solution. Another possibility is that a portion of the 13

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Ti was consumed in the reaction with oxygen since titanium oxide particles were detected.

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TEM analyses detected only M23C6 particles, and only on some of the GBs. From the pseudo-binary phase diagram (Fig. 6a), the M23C6 precipitates form in the solid state and not during solidification; therefore, they cannot act as heterogeneous nucleants for grain

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refinement. That leaves us with the possibility that Ti and C, as solutes, could produce constitutional undercooling, which could destabilize the planar solidification front. The

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partition coefficient of Ti/C < 1 (Fig. 6a), suggesting that they would be rejected into the liquid ahead of the solid-liquid interface, consistent with the EDX map of Ti (Fig. 4b) showing its enrichment in the interdendritic regions, and the presence of M23C6 precipitates (i.e., C enrichment) in the grain boundary, Fig. 6d. Solute enrichment ahead

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of the advancing solid–liquid interface could generate constitutional undercooling in the diffusion layer in front of the interface, locally reducing the nucleation barrier for

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crystallization, which results in increased number of nuclei and thus finer grains. Yamanaka [33] studied grain refinement in Co–28Cr–9W–1Si alloys by the addition of C

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(up to 0.33 wt.%) and attributed it to constitutional undercooling due to carbon rejection into the liquid ahead of the solid-liquid interface. Interestingly, although both Ti and C have individual partitioning coefficients < 1 during solidification (see pseudo-binary phase diagrams CrCoNi-Ti and CrCoNi-C in Fig. 8), our experiments have shown that adding C by itself at a level of 0.4 at.% does not result in significant grain refinement both near the mold (Fig. 9b1) as well as near the center (Fig. 9b2) whereas Ti at a level of 0.4 at.% induces partial CET near the mold surface (Fig. 9c1) but not much grain 14

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refinement in the equiaxed central region (Fig. 9c2). This suggests synergism between Ti and C probably arising from the very strong affinity of Ti for C, which slows diffusion

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ahead of the interfaces [56]. Such a slowing down of diffusion is a possible additional mechanism that could further restrict grain growth. This mechanism is different from that previously ascribed to single elements [38, 40, 57, 58], and could be a more effective

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technique for grain refinement requiring lower amounts of added elements in certain

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alloys. 4. Conclusions

(1) Microalloying additions of Ti and C induce a columnar to equiaxed transition in the shape of the as-cast grains of the CrCoNi medium-entropy alloy. The alloy with

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composition (CrCoNi)99.2Ti0.4C0.4 shows uniform equiaxed grains with a relatively fine grain size of ~75 µm. In contrast, the base CrCoNi alloy without Ti and C has long columnar grains with average lengths of ~1000 µm and widths of ~120 µm.

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(2) The yield strength, ultimate tensile strength and elongation to fracture of the

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(CrCoNi)99.2Ti0.4C0.4 MEA are all superior to those of the CoCrNi MEA at both 293 and 77 K.

(3) The columnar to equiaxed transition and grain refinement are attributed to the

segregation of Ti and C ahead of the solid-liquid interface due to their strong mutual affinity, which leads to enhanced constitutional undercooling. Acknowledgements Research sponsored by the U.S. Department of Energy, Office of Science, Basic 15

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Energy Sciences, Materials Sciences and Engineering Division, E.P.G., and the China Scholarship Council (201506165006), X.W.L. G.L. and J.P.M. acknowledge funding

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from the German Research Foundation (DFG) through project B5 and B7 of the SFB/TR 103.

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Fig. 1 (a) Schematic diagram of cast ingot showing the cross-sectional position (30 mm from bottom) where microstructures were examined. EBSD images of various as-cast

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(CrCoNi)100-2xTixCx MEAs close to the mold surface (b) and near the center (c) of the cross-section shown in (a).

Fig. 2 Grain sizes (near the center) of the MEA castings (location c in Fig. 1a) as a

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Fig. 3 EBSD images of as-cast (CrCoNi)100-2xTixCx (x=0, 0.2, 0.4 and 0.5) MEAs on

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cross-sections located 10 mm from the ingot bottom and close to the mold surface. Fig. 4 Back-scattered electron images and energy dispersive X-ray (EDX) maps of as-cast MEAs: (a) CG and (b) EG. Dendritic segregation is visible in the EDX maps. Fig. 5 Concentration profiles obtained by EDX along the black line shown in Fig. 4b.The

with red lines.

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locations where interdendrites are intersected by the concentration profile are highlighted

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Fig. 6 (a) Calculated pseudo-binary phase diagram of (CrCoNi) – Ti/C system; (b) XRD patterns of the as-cast CG and EG MEAs; (c) Representative GBs at a triple point of the

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EG MEA imaged in BSE mode showing GBs with serrations (red arrow) and smooth GBs (white arrows); (d) TEM bright-field image of serrated GB with inset showing the selected area diffraction pattern from one of the serrations. Fig. 7 (a) Engineering stress vs. engineering plastic strain; (b) True stress vs. true strain; (c) Work hardening rate (derivative of the true stress–true strain curves, dσ/dε) plotted as a function of the true stress, for CG and EG alloys. The grey area in (c) shows the necking region according to Considere’s criterion. 20

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Fig. 8 Calculated pseudo-binary phase diagrams of (a) (CrCoNi) – C and (b) (CrCoNi) – Ti systems.

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Fig. 9 Grain size and shape of CrCoNi (a) and the effects of individual additions of 0.4 at.% C (b) and 0.4 at.% Ti (c) on grain refinement.. Upper row (with subscript 1) shows EBSD images taken close to the mold interface and lower row (with subscript 2) shows

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shown in the schematic diagram in Fig. 1 (a)).

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ACCEPTED MANUSCRIPT Highlights Grain refinement of cast CrCoNi alloy is reported for the first time.




It was accomplished by microalloying with Ti and C.




Grain refinement in as-cast state enhances both tensile strength and ductility.




Mutual affinity/segregation of Ti and C induces undercooling and grain

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refinement.