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Grain boundary decohesion by nanoclustering Ni and Cr separately in CrMnFeCoNi high-entropy alloys

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Science Advances  06 Dec 2019:
Vol. 5, no. 12, eaay0639
DOI: 10.1126/sciadv.aay0639

Abstract

The loss of ductility with temperature has been widely observed in tensile tests of single-phase face-centered cubic structured high-entropy alloys (HEAs). However, the fundamental mechanism for such a ductility loss remains unknown. Here, we show that ductility loss in the CrMnFeCoNi HEA upon deformation at intermediate temperatures is correlated with cracking at grain boundaries (GBs). Nanoclustering Cr, Ni, and Mn separately at GBs, as detected by atom probe tomography, reduces GB cohesion and promotes crack initiation along GBs. We further demonstrated a GB segregation engineering strategy to avoid ductility loss by shifting the fast segregation of principal elements from GBs into preexisting Cr-rich secondary phases. We believe that GB decohesion by nanoclustering multiprincipal elements is a common phenomenon in HEAs. This study not only provides insights into understanding ductility loss but also offers a strategy for tailoring ductility-temperature relations in HEAs.

INTRODUCTION

The ductility of materials is determined primarily by strain hardening rate, which depends on alloying constituents, phase structure, microstructure, temperature, and strain rate (1). Metals become more easily deformable, accompanied by a decrease in strength and ductility, when they are heated because nucleation and motion of dislocations are thermally activated processes (2, 3). Enhancing strain rate and/or alloying appropriate elements may decrease this thermally induced softening. In addition, grain boundaries (GBs) act as barriers for the motion of dislocations, strengthening materials, and enhancing strain hardening rate. However, GBs also act as sources for nucleating the carriers of plasticity, such as dislocations and cracks. GB strengthening generally decreases with increasing deformation temperature, which would lead to the loss of ductility at intermediate temperatures if GBs are weakened by the segregation of impurity elements at GBs, and enhanced ductility at high temperatures owing to the easy motion of the GBs. For face-centered cubic (fcc) conventional metals and alloys with low stacking fault energy, such as transformation-induced plasticity and twinning-induced plasticity steels, they exhibit exceptionally high ductility due to twinning- and transformation-induced strain hardening in a wide range of temperatures (4), showing typical ductile fracture. However, recent studies have shown that single-phase fcc high- and medium-entropy alloys (HEAs/MEAs) that are composed of three or more of elements Al, Cr, Mn, Fe, Co, and Ni exhibit a rapid decrease in ductility at intermediate temperatures (400° to 800°C) (59), as summarized in Fig. 1 and fig. S1. The common feature at stress-stain curves is that these alloys exhibit strain hardening and failure shortly after the maximum tensile stress is reached. Here, we reveal the fundamental mechanism for such a ductility loss and provide guidelines for avoiding this problem.

Fig. 1 Temperature-dependent ductility.

The uniform elongation-temperature relation of various fully recrystallized single-phase fcc HEAs/MEAs with different grain sizes (59).

The ductility loss in metals and alloys can be related to the segregation of impurity elements at GBs (10, 11) and the absorption of hydrogen or helium (12), which weaken GB strength and promote crack initiation at GBs, and the precipitation of second phases (intermetallic compounds, carbides, nitrides, oxides, etc.) at GBs (13, 14), which causes high stress/strain concentration due to the incompatibility of elastic and plastic deformation between particles and matrix. In contrast to carbides and nitrides, intermetallic compound–induced embrittlement at room temperature can be relieved at high temperatures because plastic codeformation between these compounds and the matrix can reduce stress concentration and, thus, delay crack initiation (15). Taking the well-studied equiatomic Cr20Mn20Fe20Co20Ni20 HEA as an example, this alloy is a single fcc-phase solid solution after homogenization at temperatures ≥1000°C for up to 2 days and recrystallization at temperatures ≥800°C for 1 hour (5, 1618). Various intermetallic phases are precipitated primarily at GBs during annealing at intermediate temperatures for tens to hundreds of days in coarse-grained alloys (1921), while mere minutes in nanocrystalline alloys (22). Although intermetallic phases reduce the ductility of the annealed alloys at room temperature compared with the single-phase alloys, their ductility does not decrease with increasing deformation temperature (23).

During the tensile test process at intermediate temperatures in the air, intermetallic phases were not yet formed in coarse-grained Cr-Mn-Fe-Co-Ni HEAs, owing to a short testing time (a few minutes). In addition, the above-mentioned oxide and hydrogen embrittlement mechanisms can be ruled out because these alloys show good resistance to oxidation at intermediate temperatures and hydrogen embrittlement at relatively low hydrogen concentrations (24, 25). These facts suggest that temperature-induced ductility loss in fcc-phase HEAs/MEAs could be related to segregation of some of the principal elements during the short testing time, which causes GB decohesion. We examined this hypothesis in an equiatomic Cr20Mn20Fe20Co20Ni20 HEA and revealed the fast segregation of Cr, Mn, and Ni separately into GBs, which accounts for ductility loss with increasing temperature. Correspondingly, we proposed and demonstrated that preexisting precipitates can migrate the fast segregation of Cr, Mn, and Ni into the precipitates instead of at GBs, which can effectively evade ductility loss during deformation at intermediate temperatures.

RESULTS AND DISCUSSION

Ductility loss accompanied by GB cracking

Figure 2A shows the typical engineering stress-strain curves of the annealed single-phase Cr20Mn20Fe20Co20Ni20 alloy (grain size around 50 μm; see fig. S2) at seven testing temperatures. Strengths and elongations decrease with the increase in testing temperature, which falls in the results of the Cr-Mn-Fe-Co-Ni systems (Fig. 1). Pronounced serrations appear on the stress-strain curves that were obtained at 400° to 700°C. In this temperature range, the alloy exhibits fracture shortly after the ultimate tensile stress is reached. Figure 2B shows a slow decrease in the strain hardening rate with increasing temperatures from 25° to 600°C. At 700° and 800°C, the strain hardening rate decreases rapidly after yielding, with an increased necking strain.

Fig. 2 Mechanical properties of the single-phase coarse-grained Cr20Mn20Fe20Co20Ni20 HEA at different deformation temperatures.

(A) Engineering stress-strain curves and (B) strain hardening rates at 25° to 800°C.

Figure 3 compares the fracture mode of the Cr20Mn20Fe20Co20Ni20 HEA at testing temperatures of 25° and 700°C. The fracture mode changes from a fully ductile transgranular failure at 25°C (Fig. 3A) to a mixed brittle intergranular and ductile transgranular failure at 700°C (Fig. 3B). The lateral surface in the vicinity of the fracture surface at 25°C is crack free, featured with surface relief (Fig. 3C). In contrast, numerous discrete cracks are presented along GBs at the lateral surface of the sample after tensile test at 700°C, as indicated in the inserted scanning electron microscopy (SEM) image of Fig. 3D. This suggests that cracks initiate along GBs during tensile testing, leading to prefailure that accounts for the fast decrease in the ductility and strain hardening rate.

Fig. 3 Failure characteristics of the single-phase coarse-grained Cr20Mn20Fe20Co20Ni20 HEA.

(A) Fracture surface and (C) lateral surface of the originally polished tensile sample after tensile tests at 25°C. (B) Fracture surface and (D) lateral surface of the tensile sample after tensile tests at 700°C. Inset of (D), the enlarged SEM image showing many cracks along the GBs.

Nanoclustering induced GB embrittlement

Figure 4A is a transmission electron microscopy (TEM) image of the specimen tested at 700°C, showing a typical high-angle GB (marked by arrow). No second phase was detected along the GB by detailed x-ray diffraction (XRD), high-resolution TEM, electron diffraction, and SEM–EDS (energy-dispersive spectroscopy) analyses, as evidenced in fig. S3. The atom map in Fig. 4B obtained from atom probe tomography (APT) analysis shows the segregation of trace elements C and B at the GB. C and B are just impurities with concentrations of 0.2 and 0.008 atomic % (at %), respectively, in the alloys (fig. S4). The segregation of P, S, O, or other harmful elements at GBs was not detected. The isocomposition surface in Fig. 4 (C and D) and the one-dimensional (1D) compositional profiles in Fig. 4 (E and F) show substantial enrichment of Cr, Ni, and Mn in nanosized regions at the GB (around 5 nm in thickness direction). On the basis of the observation, Ni and Mn likely cosegregate at GBs, but Cr and Ni segregate separately on the GB, suggesting that they are repulsive during segregation into GBs, as shown in the normal view of the GB area in Fig. 4D.

Fig. 4 Nanosegregation at GBs.

(A) TEM image of the single-phase coarse-grained Cr20Mn20Fe20Co20Ni20 sample after tensile test at 700°C. (B) Atom map of the tip showing GB segregation of impurities C and B. Cr (24 at %) and Ni (23 at %) isocomposition surfaces (C) viewed with the GB edge-on and (D) viewed from the GB normal. 1D compositional profiles (E) along the cyan arrow and (F) along the orange arrow indicated in (C); the shaded areas represent the approximate GB regions in thickness direction.

APT data reveal that the embrittled GBs are enriched with principal elements Cr, Ni, and Mn and minimal trace elements C and B (Fig. 4, B to F). Previous works show that all principal elements in coarse-grained Cr20Mn20Fe20Co20Ni20 HEA after annealing at temperatures ≥1000°C exhibit homogeneous distribution down to the atomic scale at GBs (26, 27), without segregation except trace elements C and B. This indicates that the segregation of elements Cr, Ni, and Mn into GBs in current HEA occurred during tensile tests at 700°C. It has been shown that the residual strain generated by cold rolling can promote the diffusion of Cr into GBs during subsequent annealing processes, leading to the formation of Cr-rich σ phases in Cr-Mn-Fe-Co-Ni system HEAs (20, 28). Therefore, it is reasonable to infer that plastic deformation can facilitate the segregation of elements Cr, Ni, and Mn into GBs in current HEA. However, the low-level concentration of these elements is insufficient for precipitating Cr-rich σ, body-centered cubic (bcc) Cr, or L10-NiMn phases (table S1) (19, 20, 22). In addition, the short exposure (10 min) at 700°C is insufficient to precipitate σ phase, as confirmed by TEM (Fig. 4A). The chemical composition of the possible GB precipitates in Cr20Mn20Fe20Co20Ni20 HEA (see table S1) further confirms that Ni and Mn tend to cosegregate at GBs, while the Cr and Ni segregate separately at GBs, which is consistent with our APT result in Fig. 4. Previous studies on Fe-Mn steels have shown that the Mn segregation to GBs can induce direct embrittlement because the local magnetic contributions of Mn reduce GB cohesion (10, 29, 30). Ni and Mn cosegregated at GBs (Fig. 4F), in favor of formation of ordered L10-NiMn phase, which may decrease GB cohesion. In addition, Cr can compete for GB sites and repel Ni (Fig. 4, C to E), promoting the formation of brittle Cr-rich but Ni-poor σ phase at GBs that can reduce ductility significantly because of the weak interfaces. In addition, the segregation of trace elements C and B into GBs is beneficial to improve GB cohesion as long as M6C carbides do not precipitate (26, 31). The 1D compositional profiles in fig. S4 show that the C and B composition at Cr-enriched and Ni(Mn)-enriched GBs is very low, which is insufficient to form carbides or borides at GBs. A recent study of the single-phase fcc CoFeNi2V0.5Mo0.2 HEA demonstrated that the lack of Cr and Mn can evade ductility loss at intermediate temperatures (32). Therefore, we believe that the segregation of Cr, Ni, and Mn at GBs in Cr-Mn-Fe-Co-Ni system HEAs can embrittle GBs and, hence, lead to ductility loss with increasing deformation temperature. Note that the segregation of Cr, Ni, and Mn at GBs also occurs at relatively low temperatures (<700°C), which accounts for ductility loss at intermediate temperatures. In the current work, the sample after tensile test at 700°C, exhibiting remarkable ductility loss though maintaining a strain hardening capacity, provides a good model example to demonstrate our hypothesis: Nanoclustering induced GB embrittlement in HEAs at intermediate temperatures. We believe that GB decohesion by nanoclusters of multiprincipal elements is a common phenomenon in HEAs. Similar to Cr-rich σ phases that precipitate at intermediate temperatures (400° to 800°C) but eventually dissolve at high temperatures (>800°C) (28), nanoclusters at GBs may disappear at high temperatures; thus, there is no ductility loss. It has been shown that Cr-Mn-Fe-Co-Ni system HEAs exhibit an increased ductility and even superplasticity at temperatures over 900°C (33).

GB segregation engineering to enhance ductility

Segregation of minor alloying elements to GBs often embrittles conventional alloys (34, 35). Grain refinement is an effective way to evade GB embrittlement because more interfaces reduce the level of individual GB segregation (35). In contrast, multiple principal constitutive elements in HEAs can segregate into GBs instead of trace alloying elements. In HEAs, grain refinement can assist the segregation of principal constitutive elements to GBs and even lead to the formation of brittle intermetallic phases at GBs (20). For example, Cr-rich bcc phases and L10-NiMn phases precipitate at GBs in nanocrystalline Cr20Mn20Fe20Co20Ni20 HEA (grain size, ~50 nm) after annealing at 450°C for only 5 min, which reduces the ductility substantially (22). In addition, previous studies show that Cr20Mn20Fe20Co20Ni20 HEAs with the grain sizes from 4.4 to 155 μm all exhibit loss of ductility at intermediate temperatures, irrespective of grain size in this range (Fig. 1). Microalloying with B in Fe-Mn steels has been reported to delay temper embrittlement at 450°C (36). B can compete for GB sites and reduce GB energy and, thus, repel the embrittling element Mn due to the decreased driving force for Mn segregation (10), but in the Cr-Mn-Fe-Co-Ni system, Mn can cosegregate with Ni at GBs; thus, the minimal trace element B does not effectively compensate the embrittlement effect caused by the multiple principal constitutive element segregation. To evade the fast segregation of principal elements Ni, Cr, and Mn into GBs, we propose to migrate the fast segregation of these elements from GBs to the preexisting precipitates that are enriched with principal embrittling elements, especially for alloys used at intermediate and high temperatures, because intermetallic compounds can deform at high temperatures, although they reduce the ductility of the alloys at room temperature compared with the single-phase alloys.

On the basis of the chemical compositions of four possible GB precipitates in the Cr20Mn20Fe20Co20Ni20 HEA reported by previous studies (see table S1) (19, 20, 22), the Cr-rich σ phase absorbs Cr elements, while it repels other element segregation to GBs. Otto et al. found that GBs are rich in Ni and have no obvious segregation of embrittling elements Mn and Cr when Cr-rich σ phases were precipitated at GBs (19, 20). In addition, without Mn, Ni segregated in σ phase–free GBs is beneficial to GB cohesion. Thus, it is expected that the loss of ductility at intermediate temperatures can be evaded by preexisting precipitation of Cr-rich σ phases that can absorb embrittling Cr elements from GBs and repel other embrittling Ni, Mn element segregation to GBs. We, thus, fabricated Cr26Mn20Fe20Co20Ni14 HEAs that have a dual-phase structure (Cr-rich σ phase and fcc-phase matrix) (see fig. S5). Figure 5 shows representative engineering stress-strain curves of the Cr26Mn20Fe20Co20Ni14 HEA at the testing temperatures from 25° to 700°C. The ductility of the HEA increases with testing temperature. Although the ductility of the dual-phase HEA is lower than the single-phase HEA at the testing temperature below 500°C, the ductility of the dual-phase HEA is greater than the single-phase HEA at 600° and 700°C. No intergranular embrittlement is observed. Figure 5 (B and C) shows the lateral surface in the vicinity of the fracture after tensile tests at 500° and 700°C, respectively. Figure 5 (D and E) exhibits typical ductile dimples on fracture surfaces after tensile tests at 500° and 700°C, respectively. In contrast to single-phase samples (Fig. 3D), the dual-phase sample has very few cracks near the lateral fracture surface. This suggests that the GB segregation engineering strategy via preexisting Cr-rich precipitates is effective to evade ductility loss at intermediate temperatures in HEAs. This is because the embrittling elements—such as Cr, Mn, or Ni—can be absorbed by preexisting embrittling element–enriched precipitates at GBs, which only leads to the growth of precipitates rather than GB embrittlement. We note that the dual-phase HEA reported in this work was designed mainly with the aim of proof of the GB segregation engineering strategy to evade ductility loss by shifting the fast segregation of principal elements from GBs into preexisting embrittling element–enriched secondary phases. In our opinion, the ductility of these HEAs at intermediate temperatures can be further improved following the strategy proposed here, if the morphology, dimension, and distribution of secondary phases are optimized further.

Fig. 5 Mechanical properties and failure characteristics of the dual-phase Cr26Mn20Fe20Co20Ni14 HEA.

(A) Engineering stress-strain curves at 25° to 700°C. (B and C) Lateral fracture surfaces of the tensile samples after tensile tests at 500° and 700°C, respectively. (D and E) Fracture surfaces of the sample after tensile tests at 500° and 700°C, respectively.

CONCLUSION

We revealed the key mechanism responsible for the loss of the ductility with increasing temperature, a phenomenon widely reported in the single-phase fcc HEAs that are composed of three or more elements of Al, Cr, Mn, Fe, Co, and Ni. SEM of tensioned sample surfaces of single-phase coarse-grained Cr20Mn20Fe20Co20Ni20 HEA (grain size around 50 μm) suggested that ductility loss is correlated with the intergranular cracking. However, TEM along GBs confirmed that there are no precipitates of secondary phases along GBs. On the basis of APT analysis, we explored that ductility loss is attributed to fast nanosegregation of principal elements Ni, Cr, and Mn separately into GBs, which reduce GB cohesion and promote crack initiation along GBs. We further demonstrated a GB segregation engineering strategy to evade ductility loss by migrating fast segregation of the principal elements from GBs into preexisting Cr-rich secondary phases in an example HEA, i.e., Cr26Mn20Fe20Co20Ni14. Our work not only provides insights into understanding ductility loss but also offers a GB segregation engineering strategy of tailoring strength-ductility-temperature relations in HEAs.

MATERIALS AND METHODS

Material processing and sample preparation

The equiatomic Cr20Mn20Fe20Co20Ni20 and nonequiatomic Cr26Mn20Fe20Co20Ni14 (at %) HEAs were prepared by arc melting under an argon atmosphere using elemental ingredients with purity above 99.9 at %, subsequently homogenized in vacuum at 1200°C for 48 hours, and then cooled in the furnace. The homogenized specimens were further hot rolled at 1100°C to sheets with a thickness of 8 mm, and then cold rolled to 2.4 mm with a thickness reduction of 70%. The cold-rolled Cr20Mn20Fe20Co20Ni20 sample was subjected to annealing at 1000°C for 1 hour, followed by air cooling, leading to formation of single-phase homogeneous equiaxed grain structures with an average grain size of around 50 μm (fig. S2). The cold-rolled Cr26Mn20Fe20Co20Ni14 sample was annealed at 900°C for 1 hour to obtain a dual-phase structure, i.e., fine Cr-rich σ phases are distributed at GBs of the fcc-phase matrix with the grain size around 4 μm (fig. S5).

Mechanical and microstructural characterizations

Quasi-static tensile tests were performed on an Instron 8801 instrument at different temperatures (25° to 800°C) with an engineering strain rate of 10−3 s−1. The tensile samples were held at the desired testing temperature for ~10 min before the start of the test. Flat dog bone–shaped tensile specimens with a gauge length of 18 mm and width of 3 mm were machined from annealed samples by electric discharge machining and were ground through 1200-grit SiC paper, resulting in a final specimen thickness of 1.5 mm. The structures of the specimens were analyzed by XRD (D/MAX-2500) using Cu Kα radiation with scanning speeds of 6°/min and 0.5°/min. Microstructure characterizations were performed using an optical microscope (OM) and a JEM-2100F TEM. The nanoscale elemental distribution in the vicinity of GBs with random misorientation was characterized using APT (LEAP 3000X HR, Cameca Inc.) in laser pulsing mode with pulse rate of 250 kHz, pulse energy of 40 pJ, and specimen base temperature of 70 K. Data analysis was performed with the software IVAS 3.6.14 from Imago Scientific Instruments. The fracture surfaces of the specimens after the tensile tests were observed by SEM coupled with EDS. OM specimens were initially polished using 5000-grit SiC paper and, subsequently, carefully polished down using 1.5-μm diamond polishing compound, and etched with a solution that was composed of supersaturated copper sulfate, hydrochloric acid, and alcohol (1 mg:5 ml:5 ml). Specimens for TEM observation were thinned using mechanical grinding followed by double-jet electropolishing in a solution of 6% perchloric acid, 35% n-butyl alcohol, and 59% alcohol at −30°C and an applied voltage of 30 V. The APT tips were produced by site-specific liftout procedures from the regions including GBs using a focused ion beam instrument (FEI Helios NanoLab 600i) (37, 38).

SUPPLEMENTARY MATERIALS

Supplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/5/12/eaay0639/DC1

Fig. S1. Mechanical properties of various single-phase fcc HEAs/MEAs on the elements Al, Cr, Mn, Fe, Co, and Ni.

Fig. S2. Microstructure of the 1000°C annealed Cr20Mn20Fe20Co20Ni20 HEA.

Fig. S3. Microstructure of the 1000°C annealed Cr20Mn20Fe20Co20Ni20 sample after tensile test at 700°C.

Fig. S4. Segregation of trace elements C and B at the GB.

Fig. S5. Microstructure of the dual-phase Cr26Mn20Fe20Co20Ni14 HEA.

Table S1. Compositions of GB intermetallic phases.

This is an open-access article distributed under the terms of the Creative Commons Attribution-NonCommercial license, which permits use, distribution, and reproduction in any medium, so long as the resultant use is not for commercial advantage and provided the original work is properly cited.

REFERENCES AND NOTES

Acknowledgments: Funding: This work is supported by the Science Fund for Creative Research Groups (61271043) and the Research Council at the University of Nebraska–Lincoln. The research was performed, in part, in the Nebraska Nanoscale Facility: National Nanotechnology Coordinated Infrastructure and the Nebraska Center for Materials and Nanoscience, which are supported by the National Science Foundation under Award ECCS 1542182 and the Nebraska Research Initiative. Author contributions: K.M., X.B., and J.W. designed the study and wrote the first draft of this manuscript. K.M. prepared the alloys and performed mechanical tests and TEM observations. L.L. and Z.L. conducted the APT experiments and interpreted the results. All the authors analyzed and interpreted the data and commented on the manuscript. Competing interests: The authors declare that they have no competing interests. Data and materials availability: All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials. Additional data related to this paper may be requested from the authors.
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