- Email: [email protected]
Understanding on toughening mechanism of bioinspired bulk metallic glassy composites by thermal spray additive manufacturing
Scripta Materialia 177 (2020) 112–117
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Understanding on toughening mechanism of bioinspired bulk metallic glassy composites by thermal spray additive manufacturing Cheng Zhang a,1, Wei Wang b,1, Wei Xing a, Lin Liu a,∗ a
School of Materials Science and Engineering and State Key Lab for Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, Wuhan 430074, China b National Engineering Laboratory for Modern Materials Surface Engineering Technology, The Key Lab of Guangdong for Modern Surface Engineering Technology, Guangdong Institute of New Materials, Guangdong 510651, China
a r t i c l e
i n f o
Article history: Received 4 October 2019 Accepted 11 October 2019
Keywords: Bulk metallic glasses composite Thermal spray additive manufacturing Interfaces Mechanical property Toughening mechanism
a b s t r a c t Toughening mechanism of bioinspired bulk metallic glassy composites (BMGCs) reinforced with various volume fractions of ductile stainless-steel (SS) phase fabricated via newly developed thermal spray additive manufacturing was investigated. Microstructure characterizations and nanoindentation suggested a strong interface-bonding along with nanoscale compositional gradient formed between the metallic glass and SS phases. Finite element modelling and theoretical analysis revealed that the increase in toughness arose from the nacre-like soft/hard phase configuration, which exhibits the effective stress-shielding effect against crack propagation, thus mitigating stress concentration and blocking crack propagation. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Bulk metallic glasses (BMGs), especially Fe-based BMGs, are attracting considerable attention in last decades owing to their unique properties, including high strength and unique functions [1–8]. However, the size limitation, poor manufacturability, and low ductility are the major hurdles restraining the applications of Fe-based BMGs as structural materials [7]. To surmount the limitation of glass-formation ability (GFA), the approach of processing based on powder metallurgy was developed. For example, spark plasma sintering (SPS) technique had been employed to synthesize large-sized Fe-based BMGs. However, the applied stress needed to obtain fully densified parts exceeds 600 MPa, thus the expensive WC mould are required, making this approach non-applicable in practical applications [9,10]. Additive manufacturing (AM) technology has been recently utilized to fabricate large-sized BMG parts with complex geometries [11–18]. Up to now, some achievements have been realized for Zr-, Ti-, Fe-, and Al-based BMGs using laser-based AM processes such as selective laser melting-SLM [11,12,15–18]. However, the Fe-based BMG parts printed with SLM usually show unexpectedly poor mechanical properties due to the presence of defects in the as-printed parts, such as micro-cracks, pores, and partial crystallization in heat affect zones [11,16,18]. Recently, we developed a novel thermal spray 3D printing (TS3DP) technique to fabricate large Fe-based BMG and BMG composites (BMGCs), in which ∗ 1
Corresponding author. These authors contributed equally to this work. E-mail address: lliu20 0 [email protected] (L. Liu).
https://doi.org/10.1016/j.scriptamat.2019.10.017 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the semi-molten particles solidified with a high cooling rate and ultimately assembled into a bioinspired nacre-like structure [19]. With this technique, the as-printed BMG (with the composition of Fe48 Mo14 Cr15 Y2 C15 B6 ) exhibited an enhanced fracture toughness of KJ =∼13 MPa m1/2 , which is two times higher than the fracture toughness of the as-cast counterpart. By further modifying the alloy composition using mixtures of metallic glass (MG) and ductile stainless-steel (SS) powders, a higher toughness of ∼21 MPa m1/2 was achieved in the resultant TS3DP BMGC with 50 wt% SS [19]. However, the detailed mechanisms for the enhancement of fracture toughness in the SS reinforced BMG composites are still not well understood, this issue is essentially important for the development of 3D-printed BMG composites with excellent mechanical properties. In nacre-like composites, the interfacial structure of brick/mortar and the relative content of mortar usually play a vital role in mechanical properties of the composites [20]. Therefore, there is a compelling need to understand the correlation between interfacial structure and mechanical properties of the SS reinforced BMGCs. This issue is important for the design and optimization of high-performance BMGCs by TS3DP. In the current work, Fe-based BMGCs reinforced with two different volume fractions of ductile SS phase prepared by the TS3DP method was characterized by transmission electron microscopy (TEM) and nanoindentation. The toughening mechanisms are discussed in details on the basis of finite element simulation and theoretic analysis.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
113
Fig. 1. (a) EBSD map showing the distribution and orientation of the SS phase in the amorphous matrix; (b) TEM image, (c) HRTEM image, and (d) HAADF-STEM image, showing the interfacial structure between MG and SS phases. (e) EDX scanning across the MG/SS interface. (f) Load-depth curves of the amorphous and SS phases; (g) load-depth curves measured from the MG-SS interface with different applying loads. The inset in (g) shows an optical micrograph of an indent at the interface region.
Amorphous Fe48 Mo14 Cr15 Y2 C15 B6 (at.%) powder was prepared through high-pressure Ar gas atomization, and particles in the range of 33–55 μm were sieved out for thermal spray. The MG particles and commercial 316 L SS powders (30–60 μm) in the mass fraction of 20%, and 50% were mixed by ball milling. The BMGCs were fabricated via TS3DP technique using the same spraying parameters as described in our previous work [19]. The crystallographic orientation of the SS phase in the printed samples was examined via electron back scattered diffraction (EBSD). The interfacial structure between the MG splat and SS splat in the BMGC was examined by transmission electron microscopy (TEM, FEI Tecnai G2 F20) coupled with an energy dispersive X-ray spectroscopy (EDX). TEM samples were prepared by ion milling (Gatan 691). A TTX-NHT3 nanoindentation tester was employed to measure the hardness and Young’s moduli of single MG and SS splat of the BMGC with an applied load of 8 mN. The bonding strength between MG and SS phases was further evaluated by nanoindentation at contact loads between 100 and 400 mN. Strength and plasticity of the BMGC were measured in uniaxial compression tests at a strain rate of 10−4 s−1 . The dimensions of the samples used for compression testing were 3 mm × 6 mm. Experiments to evaluate fracture toughness were conducted with single-edged pre-notched samples having dimensions of 1.5 × 3 × 20 mm3 . Notches with root radius of 144 μm and depth of 750 μm were carefully produced by low-speed wire cutting. ABAQUS software was used to simulate the stress distribution in the samples upon three-point bending, and the details of the finite element modelling (FEM) are described in the Supplementary Materials. Two nearly fully densified BMGC components containing 20% and 50 wt.% ductile SS phases (named as S20 and S50 thereafter) were prepared with the TS3DP technique. EBSD analysis (Fig. 1a) indicated that the austenite SS phase (colored regions) had a random orientation, and was homogenously distributed in the amorphous matrix (the remaining dark regions). TEM examinations
were performed with the S50 specimen to examine the interfacial structure between MG and SS phases. A bright-field TEM image of an interfacial region is shown in Fig. 1b, which indicates clearly a perfect metallurgical bonding between MG and SS phases. A high density of dislocations in the SS phase indicates that the SS particles underwent severe plastic deformation during thermal spraying. Since the re-heating temperature on the previously deposited layers was normally below 473 K [21], far below the recrystallization temperature of the SS phase (> 873 K [22]), thus SS recrystallization could not occur during thermal spray process. Fig. 1c shows a HR-TEM image of the interface. Regular lattice fringes were visible in the SS phase, which were clearly distinct from the maze-like patterns of the amorphous phase. The corresponding fast Fourier transform patterns of the two phases indicate that the crystalline (FCC) and amorphous structures of the two phases remains unaffected. In addition, the high-angle annular dark-field (HAADF)-STEM imaging coupled with EDX analysis of the interface (Fig. 1d and e) verified the presence of a compositional gradient region of ∼45 nm in width, which indicates that atomic diffusion occurred between the two phases in TS3DP process. Nanoindentation test was performed to probe the hardness of the single-splat amorphous and SS phases (Fig. 1f), which revealed that the hardness of the amorphous phase and SS phase to be 19.6 ± 0.9 and 5.9 ± 0.5 GPa, respectively. To investigate the bonding strength of the MG/SS interface, nanoindentation tests were intentionally carried out at the interfacial region with a high loading force up to 400 mN. No pop-in or step events corresponding to interface debonding was observed in the corresponding loaddepth curves (Fig. 1g), and no cracks were observed at the interface (Fig. 1g, inset). These results demonstrate that the interfacial bonding between the amorphous and SS phases is strong and tough enough, and can withstand high contact stress and severe plastic deformation. Fig. 2a shows the stress-strain curves of the two BMGC samples (S20 and S50) under compression. To assess the anisotropic
114
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
Fig. 2. (a) Compressive stress–strain curves of the samples measured in two directions, wherein the load was applied parallelly to the building direction (direction A) or perpendicularly to the building direction (direction B). Fracture micrographs of sample S50 after compressive tests (b, c) in direction A and (d, e) in direction B. (f) Comparison of the fracture toughness of our bioinspired BMG/SS composites with other nacre-like bioinspired hybrid materials [25–29]. Insets show the nacre-like microstructure in three selected composites including BMG/SS and Al2 O3 /Epoxy by thermal spraying, and Al2 O3 /BMG by freeze casting. (g) Fracture strength vs. fracture toughness plot of Febased BMG and BMGCs processed by various techniques. Several BMGC mechanical parts with complex geometries, such as fans and screwdrivers (insets) were successfully manufactured via TS3DP followed by machining.
properties, the samples being perpendicular to and parallel to the building direction (i.e., directions A and B, respectively) were tested by compression. The fracture strength (σ f ) in direction A (perpendicular to the lamellae) is clearly higher than that in direction B (parallel to the lamellae) in both composites, as what were commonly observed in nacre and bioinspired nacre-like materials. However, the difference in strength between in direction A and in direction B decreases remarkably as the SS mass fraction increases, suggesting that anisotropy could mitigate with high content of SS phase. In addition, it was also found that the added SS phase improved the plasticity of the BMG, especially in direction A, as indicated by plastic strain of 2% achieved for S50, 3 times higher than that for S20 (0.7%). Fractographic analysis was performed to understand the fracturing mechanisms associated with the anisotropy. Herein, the S50 specimen was selected as an example. Fig. 2b shows a typical shear morphology of the fractured surface in S50 loaded along direction A. The shear angle of approximately 41° was located, which is similar to that in the “plastic” Febased BMGCs prepared by copper mould casting [23]. This fracture behavior indicates that the Fe-based BMGCs prepared by TS3DP in
this work may have good toughness. The dimpling fracture morphology and the corresponding EDX result (Fig. 2c, inset) further confirm a ductile fracture behavior in the SS phase, although the MG phase still shows brittle feature. Splitting fracture in the sample loaded along direction B is seen in Fig. 2d. To clarify the issue that how cracks interacted with the SS phase, the sample S50 (along direction B) was further subjected to a compression test, in which a part of the side surfaces of the rod was first polished to a mirror finish, then compressed to fracture. Two large longitudinal cracks propagated mainly along the intersplats and were nearly parallel to the loading axis (Fig. 2e). In the region without SS phase, cracks behave similar to TS3DP BMG, i.e., cracks initiate at some weak sites and then propagate along the weak lamellar interfaces. However, once cracks encounter the ductile SS phase, they either cut through the SS phase or divide into multiple microcracks (Fig. S2), and then propagate again along the lamellar interfaces until meeting another SS phase. This crack propagation behavior also agrees with the viewpoint of fracture mechanics for composites with hard/soft phases, which will be discussed in the later section.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
115
Fig. 3. Contour plots of simulated third stress in (a) monolithic BMG; (b) S20 and (c) S50 at the same displacement during three-point bending. (d) The calculated stress distribution along the center line of the notch (in parallel direction). (e) Crack driving force (Jtip ) as a function of distance (L1 ) against the first interface between the SS and amorphous phase, showing how a long crack interacts within the BMGC which contains a single and soft-SS interlayer between two amorphous splats.
The real fracture toughness was measured in three-point bending tests using specimens notched in the direction A. Because of the nonlinearity in load-displacement curves of TS3DP BMGCs (not shown here), we employed the J-integral characterization, which is practicable to describe nonlinear-elastic fracture behavior of materials, to calculate the fracture toughness (KJ ) of the BMGCs by following Eqs. (1) and (2) [24]:
KJ = J=
J E 0 . 5 1 − ν2
1.9 A B (W − a )
(1)
(2)
where E is Young’s modulus, ν is Poisson’s ratio (ν = 0.3), J and KJ are the nonlinear strain-energy release rate and fracture toughness, respectively, A is the total area under the load-displacement curve. B, W and a are the thickness, width and notch depth of the specimens, respectively. KJ of S20 and S50 are calculated to be approximately 16 and 21 MPa m1/2 , respectively. These values far exceed the toughness of monolithic BMG processed by either TS3DP (13 MPa m1/2 ) or casting (5 MPa m1/2 ) [19], and is equivalent to the best nacre-like materials developed previously [25–29] (see Fig. 2f). The fracture toughness of the samples increased as the mass fraction of the SS phase increased, which verified the toughening effect of the SS phase. To exhibit the advance of the TS3DP technique used in this work, we compare the fracture strength and
toughness of the TS3DP-BMGCs with those of other BMGs with similar compositions and those Fe-based BMGCs processed by SLM, SPS and casting (Table S1), as shown in Fig. 2g. Generally, casting and SPS-processed BMGs have high fracture strength but low fracture toughness. The BMGCs with ductile phases (i.e., Cu and CuNi alloy) prepared by SLM have relatively good toughness, but exhibit low strength. Our BMGCs prepared by TS3DP appear at the upper-right regime of the plot, highlighting a good combination of strength and toughness. To essentially uncover the toughening mechanism, the stress fields in the pre-notched BMG composites under three-point bending were visualized in FEM simulation, in comparison with that in monolithic BMG. The contour plots of the third invariant of the stress tensor (determining the volume changes in deformation) in the specimens with equal bending displacement are shown in Fig. 3a-c. The stress level was displayed as a function of the distance from the notch tip (Fig. 3d). A large gradient from tensile stress to compressive stress was observed in the monolithic BMG upon bending. However, as the SS fraction increases, fluctuations in the stress invariant are remarkably reduced within the distance of 350 μm from the notch tip. This would lead to reduction of the driving force of crack propagation. The distribution of Mises stress (Fig. S3) also verified the shielding effect of the SS phase since the stress concentration is effectively mitigated and the stress is widely distributed in the S20 and S50. To clarify the influence of the SS phase on crack propagation, we performed theoretical calculation
116
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
cal calculation revealed that increase in toughness arose from the unique laminar soft/hard phase configuration which exhibited the effective stress-shielding effect against crack propagation. This approach may provide new options for the design and manufacturing of strong, tough BMGCs for structural applications. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments
Fig. 4. Compressive stress–strain curve of a Fe48 Mo14 Cr15 Y2 C15 B6 BMG composite/FeCoCrNiMn high-entropy alloy (HEA) sample (Ø 2 × 4 mm) prepared via TS3DP technique (the volume fraction of HEA is 80%). Inset is the XRD pattern of the composite, showing a dual structure consisting of both amorphous phase and fcc solid solution.
This work was financially supported by the National Key R&D Program of China (No. 2016YFB1100101), National Natural Science Foundation of China (Grant nos. 51531003, 51771077), the Technical Innovation Project of Hubei Province (2017AAA018) and GDAS’s Project of Science Technology Development (2018GDASCX-0402). The authors are grateful to the Analytical and Testing Center in HUST for technical assistance. Supplementary materials
of the driving force for crack growth in the TS3DP-BMGC according to the inhomogeneous model proposed by Sistaninia and Kolednik [30] (see the details in the Supplementary Materials). A general requirement for crack propagation is that the driving force in front of the crack tip (Jtip ) that exceeds the intrinsic fracture resistance of the material. The calculated Jtip of the crack propagating across the amorphous/SS interface is shown in Fig. 3e. It is seen that Jtip increased sharply as the crack approached interface I, implying that crack would easily penetrate the interface and enter the soft-SS interlayer. Nevertheless, Jtip rapidly decreased to a level far below the far-field J-integral in the amorphous matrix (Jfar ) once the crack entered the SS layer due to the shielding effect. Jtip reached its lowest value at the interface II. In this case, crack tended to deflect and propagate along the intersplat, where the intrinsic resistance to crack propagation was low. To break the next hard-amorphous layer, extra energy is required, thus crack resistance rises steeply. This result indicates that alternative softto-hard materials arrangement plays an essential role in toughening the BMGCs. Similar toughening effect was also achieved in a MG/amorphous-carbon hybrid materials [31]. To present the applicability of the TS3DP technique, we have manufactured several mechanical parts, such as fans and screwdrivers (Fig. 3g, insets). These examples demonstrate that the TS3DP technique is able to fabricate large-sized BMGC component with complex geometries. Moreover, the strategy of preparing nacre-like structure with hard and soft phases using TS3DP can also be adopted to other BMG systems or adding other ductile reinforcements to producing bioinspired hybrid architectures. For instance, we also fabricated a BMG/high-entropy alloy (HEA) dual-structured composite via TS3DP. The fcc FeCoCrNiMn HEA was chosen here owing to its good ductility and damage tolerance [32]. The resultant BMG/HEA composite exhibited a high-strength of > 1.5 GPa and reasonable plasticity of ∼3% (Fig. 4). Hence, this approach may provide new options for design and manufacturing of strong, toughening BMG composites for structural applications. In summary, a strong interfacial bonding with nanoscale compositional gradient between amorphous and SS phases was observed in bioinspired bulk metallic glassy composites prepared by thermal spray additive manufacturing technique. We demonstrated that the soft-SS reinforcement significantly enhanced the mechanical properties of Fe-based BMG. FEM simulations and theoreti-
Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.10. 017. References [1] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [2] A. Inoue, A. Takeuchi, Acta Mater. 59 (2011) 2243–2267. [3] Z. Liu, W. Chen, J. Carstensen, J. Carstensen, J. Ketkaew, R.M.O. Mota, J.K. Guest, J. Schroers, Acta Mater. 105 (2016) 35–43. [4] C. Suryanarayana, A. Inoue, Int. Mater. Rev. 58 (2013) 131–166. [5] M.J. Duarte, J. Klemm, S.O. Klemm, K.J.J. Mayrhofer, M. Stratmann, S. Borodin, A.H. Romero, M. Madinehei, D. Crespo, J. Serrano, S.S.A. Gerstl, P.P. Choi, D. Raabe, F.U. Renner, Science 341 (2013) 372–376. [6] C. Zhang, H. Zhou, L. Liu, Acta Mater. 72 (2014) 239–251. [7] H.X. Li, Z.C. Lu, S.L. Wang, Y. Wu, Z.P. Lu, Prog. Mater. Sci. 103 (2019) 235–318. [8] L.C. Zhang, Z. Jia, F. Lu, S.X. Liang, J. Lu, Prog. Mater. Sci. 105 (2019) 100576. [9] A. Singh, S.R. Paital, A. Andapally, N.B. Dahotre, S.P. Harimkar, Adv. Eng. Mater. 14 (2012) 400–407. [10] J.P. Kelly, S.M. Fuller, K. Seo, E. Novitskaya, V. Eliasson, A.M. Hodge, O.A. Graeve, Mater. Des. 93 (2016) 26–38. [11] S. Pauly, L. Lober, R. Petters, M. Stoica, S. Scudino, U. Kuhn, J. Eckert, Mater. Today 16 (2013) 37–41. [12] X.P. Li, M.P. Roberts, S.O. Keeffe, T.B. Sercombe, Mater. Des. 112 (2016) 217–226. [13] Y.Y. Shen, Y.Q. Li, C. Chen, H.L. Tsai, Mater. Des. 117 (2017) 213–222. [14] M.A. Gibson, N.M. Mykulowycz, J. Shim, R. Fontana, P. Schmitt, A. Roberts, J. Ketkaew, L. Shao, W. Chen, P. Bordeeniothikasem, J.S. Myerberg, R. Fulop, M.D. Verminski, E.M. Sachs, Y.M. Chiang, C.A. Schuh, A.J. Hart, J. Schroers, Mater. Today 21 (2018) 697–702. [15] C. Zhang, X.M. Li, S.Q. Liu, H. Liu, L.J. Yu, L. Liu, J. Alloys Compd. 790 (2019) 963–973. [16] Z. Mahhooba, L. Thorsson, M. Unosson, P. Skoglund, H. West, T. Horn, C. Rock, E. Vogli, O. Harrysson, Appl. Mater. Today 11 (2018) 1–6. [17] C. Yang, C. Zhang, L. Liu, J. Mater. Chem. A 6 (2018) 20992–21002. [18] N. Li, J.J. Zhang, W. Xing, Di. Ouyang, L. Liu, Mater. Des. 143 (2018) 285–296. [19] C. Zhang, W. Wang, Y.C. Li, Y.G. Yang, Y. Wu, L. Liu, J. Mater. Chem. A 6 (2018) 6800–6805. [20] R.P. Wilkerson, B. Gludovatz, J. Watts, A.P. Tomsia, G.E. Hilmas, R.O. Ritchie, Acta Mater. 148 (2017) 147–155. [21] P. Bansal, P.H. Shipway, S.B. Leen, Acta Mater. 55 (2007) 5089–5101. [22] S.J. Wang, T. Jozaghi, I. Karaman, R. Arroyave, Y.I. Chumlyakov, Mater. Sci. Eng. A 694 (2017) 121–131. [23] S.F. Guo, L. Liu, N. Li, Y. Li, Scr. Mater. 62 (2010) 329–332. [24] M.E. Launey, E. Munch, D. Alsem, H.B. Barth, E. Saiz, A.P. Tomsia, R.O. Ritchie, Acta Mater. 57 (2009) 2919–2932. [25] E. Munch, M.E. Launey, D.H. Alsem, E. Saiz, A.P. Tomsia, R.O. Ritchie, Science 322 (2018) 1516–1520. [26] R.P. Wilkerson, B. Gludovatz, J. Watts, A.P. Tomsia, G.E. Hilmas, R.O. Ritchie, Adv. Mater. 28 (2016) 10061–10067. [27] A. Wat, J.I. Lee, C.W. Ryu, B. Gludovatz, J. Kim, A.P. Tomsia, T. Ishikawa, J. Schmitz, A. Meyer, M. Alfreider, D. Kiener, E.S. Park, R.O. Ritchie, Nat. Commun. 10 (2019) 961.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117 [28] G. Dwivedi, K. Flynn, M. Resnick, S. Sampath, A. Gouldstone, Adv. Mater. 27 (2015) 3073–3078. [29] F. Bouville, E. Maire, S. Meille, V.D. Moortele B, A.J. Stevenson, S. Deville, Nat. Mater. 13 (2014) 508–514. [30] M. Sistaninia, O. Kolednik, Eng. Fract. Mech. 130 (2014) 21–41.
117
[31] W. Wang, C. Zhang, Z.W. Zhang, Y.C. Li, M. Yasir, H.T. Wang, L. Liu, Sci. Rep. 7 (2017) 4084. [32] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, Science 345 (2014) 1153–1158.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Understanding on toughening mechanism of bioinspired bulk metallic glassy composites by thermal spray additive manufacturing Cheng Zhang a,1, Wei Wang b,1, Wei Xing a, Lin Liu a,∗ a
School of Materials Science and Engineering and State Key Lab for Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, Wuhan 430074, China b National Engineering Laboratory for Modern Materials Surface Engineering Technology, The Key Lab of Guangdong for Modern Surface Engineering Technology, Guangdong Institute of New Materials, Guangdong 510651, China
a r t i c l e
i n f o
Article history: Received 4 October 2019 Accepted 11 October 2019
Keywords: Bulk metallic glasses composite Thermal spray additive manufacturing Interfaces Mechanical property Toughening mechanism
a b s t r a c t Toughening mechanism of bioinspired bulk metallic glassy composites (BMGCs) reinforced with various volume fractions of ductile stainless-steel (SS) phase fabricated via newly developed thermal spray additive manufacturing was investigated. Microstructure characterizations and nanoindentation suggested a strong interface-bonding along with nanoscale compositional gradient formed between the metallic glass and SS phases. Finite element modelling and theoretical analysis revealed that the increase in toughness arose from the nacre-like soft/hard phase configuration, which exhibits the effective stress-shielding effect against crack propagation, thus mitigating stress concentration and blocking crack propagation. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Bulk metallic glasses (BMGs), especially Fe-based BMGs, are attracting considerable attention in last decades owing to their unique properties, including high strength and unique functions [1–8]. However, the size limitation, poor manufacturability, and low ductility are the major hurdles restraining the applications of Fe-based BMGs as structural materials [7]. To surmount the limitation of glass-formation ability (GFA), the approach of processing based on powder metallurgy was developed. For example, spark plasma sintering (SPS) technique had been employed to synthesize large-sized Fe-based BMGs. However, the applied stress needed to obtain fully densified parts exceeds 600 MPa, thus the expensive WC mould are required, making this approach non-applicable in practical applications [9,10]. Additive manufacturing (AM) technology has been recently utilized to fabricate large-sized BMG parts with complex geometries [11–18]. Up to now, some achievements have been realized for Zr-, Ti-, Fe-, and Al-based BMGs using laser-based AM processes such as selective laser melting-SLM [11,12,15–18]. However, the Fe-based BMG parts printed with SLM usually show unexpectedly poor mechanical properties due to the presence of defects in the as-printed parts, such as micro-cracks, pores, and partial crystallization in heat affect zones [11,16,18]. Recently, we developed a novel thermal spray 3D printing (TS3DP) technique to fabricate large Fe-based BMG and BMG composites (BMGCs), in which ∗ 1
Corresponding author. These authors contributed equally to this work. E-mail address: lliu20 0 [email protected] (L. Liu).
https://doi.org/10.1016/j.scriptamat.2019.10.017 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the semi-molten particles solidified with a high cooling rate and ultimately assembled into a bioinspired nacre-like structure [19]. With this technique, the as-printed BMG (with the composition of Fe48 Mo14 Cr15 Y2 C15 B6 ) exhibited an enhanced fracture toughness of KJ =∼13 MPa m1/2 , which is two times higher than the fracture toughness of the as-cast counterpart. By further modifying the alloy composition using mixtures of metallic glass (MG) and ductile stainless-steel (SS) powders, a higher toughness of ∼21 MPa m1/2 was achieved in the resultant TS3DP BMGC with 50 wt% SS [19]. However, the detailed mechanisms for the enhancement of fracture toughness in the SS reinforced BMG composites are still not well understood, this issue is essentially important for the development of 3D-printed BMG composites with excellent mechanical properties. In nacre-like composites, the interfacial structure of brick/mortar and the relative content of mortar usually play a vital role in mechanical properties of the composites [20]. Therefore, there is a compelling need to understand the correlation between interfacial structure and mechanical properties of the SS reinforced BMGCs. This issue is important for the design and optimization of high-performance BMGCs by TS3DP. In the current work, Fe-based BMGCs reinforced with two different volume fractions of ductile SS phase prepared by the TS3DP method was characterized by transmission electron microscopy (TEM) and nanoindentation. The toughening mechanisms are discussed in details on the basis of finite element simulation and theoretic analysis.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
113
Fig. 1. (a) EBSD map showing the distribution and orientation of the SS phase in the amorphous matrix; (b) TEM image, (c) HRTEM image, and (d) HAADF-STEM image, showing the interfacial structure between MG and SS phases. (e) EDX scanning across the MG/SS interface. (f) Load-depth curves of the amorphous and SS phases; (g) load-depth curves measured from the MG-SS interface with different applying loads. The inset in (g) shows an optical micrograph of an indent at the interface region.
Amorphous Fe48 Mo14 Cr15 Y2 C15 B6 (at.%) powder was prepared through high-pressure Ar gas atomization, and particles in the range of 33–55 μm were sieved out for thermal spray. The MG particles and commercial 316 L SS powders (30–60 μm) in the mass fraction of 20%, and 50% were mixed by ball milling. The BMGCs were fabricated via TS3DP technique using the same spraying parameters as described in our previous work [19]. The crystallographic orientation of the SS phase in the printed samples was examined via electron back scattered diffraction (EBSD). The interfacial structure between the MG splat and SS splat in the BMGC was examined by transmission electron microscopy (TEM, FEI Tecnai G2 F20) coupled with an energy dispersive X-ray spectroscopy (EDX). TEM samples were prepared by ion milling (Gatan 691). A TTX-NHT3 nanoindentation tester was employed to measure the hardness and Young’s moduli of single MG and SS splat of the BMGC with an applied load of 8 mN. The bonding strength between MG and SS phases was further evaluated by nanoindentation at contact loads between 100 and 400 mN. Strength and plasticity of the BMGC were measured in uniaxial compression tests at a strain rate of 10−4 s−1 . The dimensions of the samples used for compression testing were 3 mm × 6 mm. Experiments to evaluate fracture toughness were conducted with single-edged pre-notched samples having dimensions of 1.5 × 3 × 20 mm3 . Notches with root radius of 144 μm and depth of 750 μm were carefully produced by low-speed wire cutting. ABAQUS software was used to simulate the stress distribution in the samples upon three-point bending, and the details of the finite element modelling (FEM) are described in the Supplementary Materials. Two nearly fully densified BMGC components containing 20% and 50 wt.% ductile SS phases (named as S20 and S50 thereafter) were prepared with the TS3DP technique. EBSD analysis (Fig. 1a) indicated that the austenite SS phase (colored regions) had a random orientation, and was homogenously distributed in the amorphous matrix (the remaining dark regions). TEM examinations
were performed with the S50 specimen to examine the interfacial structure between MG and SS phases. A bright-field TEM image of an interfacial region is shown in Fig. 1b, which indicates clearly a perfect metallurgical bonding between MG and SS phases. A high density of dislocations in the SS phase indicates that the SS particles underwent severe plastic deformation during thermal spraying. Since the re-heating temperature on the previously deposited layers was normally below 473 K [21], far below the recrystallization temperature of the SS phase (> 873 K [22]), thus SS recrystallization could not occur during thermal spray process. Fig. 1c shows a HR-TEM image of the interface. Regular lattice fringes were visible in the SS phase, which were clearly distinct from the maze-like patterns of the amorphous phase. The corresponding fast Fourier transform patterns of the two phases indicate that the crystalline (FCC) and amorphous structures of the two phases remains unaffected. In addition, the high-angle annular dark-field (HAADF)-STEM imaging coupled with EDX analysis of the interface (Fig. 1d and e) verified the presence of a compositional gradient region of ∼45 nm in width, which indicates that atomic diffusion occurred between the two phases in TS3DP process. Nanoindentation test was performed to probe the hardness of the single-splat amorphous and SS phases (Fig. 1f), which revealed that the hardness of the amorphous phase and SS phase to be 19.6 ± 0.9 and 5.9 ± 0.5 GPa, respectively. To investigate the bonding strength of the MG/SS interface, nanoindentation tests were intentionally carried out at the interfacial region with a high loading force up to 400 mN. No pop-in or step events corresponding to interface debonding was observed in the corresponding loaddepth curves (Fig. 1g), and no cracks were observed at the interface (Fig. 1g, inset). These results demonstrate that the interfacial bonding between the amorphous and SS phases is strong and tough enough, and can withstand high contact stress and severe plastic deformation. Fig. 2a shows the stress-strain curves of the two BMGC samples (S20 and S50) under compression. To assess the anisotropic
114
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
Fig. 2. (a) Compressive stress–strain curves of the samples measured in two directions, wherein the load was applied parallelly to the building direction (direction A) or perpendicularly to the building direction (direction B). Fracture micrographs of sample S50 after compressive tests (b, c) in direction A and (d, e) in direction B. (f) Comparison of the fracture toughness of our bioinspired BMG/SS composites with other nacre-like bioinspired hybrid materials [25–29]. Insets show the nacre-like microstructure in three selected composites including BMG/SS and Al2 O3 /Epoxy by thermal spraying, and Al2 O3 /BMG by freeze casting. (g) Fracture strength vs. fracture toughness plot of Febased BMG and BMGCs processed by various techniques. Several BMGC mechanical parts with complex geometries, such as fans and screwdrivers (insets) were successfully manufactured via TS3DP followed by machining.
properties, the samples being perpendicular to and parallel to the building direction (i.e., directions A and B, respectively) were tested by compression. The fracture strength (σ f ) in direction A (perpendicular to the lamellae) is clearly higher than that in direction B (parallel to the lamellae) in both composites, as what were commonly observed in nacre and bioinspired nacre-like materials. However, the difference in strength between in direction A and in direction B decreases remarkably as the SS mass fraction increases, suggesting that anisotropy could mitigate with high content of SS phase. In addition, it was also found that the added SS phase improved the plasticity of the BMG, especially in direction A, as indicated by plastic strain of 2% achieved for S50, 3 times higher than that for S20 (0.7%). Fractographic analysis was performed to understand the fracturing mechanisms associated with the anisotropy. Herein, the S50 specimen was selected as an example. Fig. 2b shows a typical shear morphology of the fractured surface in S50 loaded along direction A. The shear angle of approximately 41° was located, which is similar to that in the “plastic” Febased BMGCs prepared by copper mould casting [23]. This fracture behavior indicates that the Fe-based BMGCs prepared by TS3DP in
this work may have good toughness. The dimpling fracture morphology and the corresponding EDX result (Fig. 2c, inset) further confirm a ductile fracture behavior in the SS phase, although the MG phase still shows brittle feature. Splitting fracture in the sample loaded along direction B is seen in Fig. 2d. To clarify the issue that how cracks interacted with the SS phase, the sample S50 (along direction B) was further subjected to a compression test, in which a part of the side surfaces of the rod was first polished to a mirror finish, then compressed to fracture. Two large longitudinal cracks propagated mainly along the intersplats and were nearly parallel to the loading axis (Fig. 2e). In the region without SS phase, cracks behave similar to TS3DP BMG, i.e., cracks initiate at some weak sites and then propagate along the weak lamellar interfaces. However, once cracks encounter the ductile SS phase, they either cut through the SS phase or divide into multiple microcracks (Fig. S2), and then propagate again along the lamellar interfaces until meeting another SS phase. This crack propagation behavior also agrees with the viewpoint of fracture mechanics for composites with hard/soft phases, which will be discussed in the later section.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
115
Fig. 3. Contour plots of simulated third stress in (a) monolithic BMG; (b) S20 and (c) S50 at the same displacement during three-point bending. (d) The calculated stress distribution along the center line of the notch (in parallel direction). (e) Crack driving force (Jtip ) as a function of distance (L1 ) against the first interface between the SS and amorphous phase, showing how a long crack interacts within the BMGC which contains a single and soft-SS interlayer between two amorphous splats.
The real fracture toughness was measured in three-point bending tests using specimens notched in the direction A. Because of the nonlinearity in load-displacement curves of TS3DP BMGCs (not shown here), we employed the J-integral characterization, which is practicable to describe nonlinear-elastic fracture behavior of materials, to calculate the fracture toughness (KJ ) of the BMGCs by following Eqs. (1) and (2) [24]:
KJ = J=
J E 0 . 5 1 − ν2
1.9 A B (W − a )
(1)
(2)
where E is Young’s modulus, ν is Poisson’s ratio (ν = 0.3), J and KJ are the nonlinear strain-energy release rate and fracture toughness, respectively, A is the total area under the load-displacement curve. B, W and a are the thickness, width and notch depth of the specimens, respectively. KJ of S20 and S50 are calculated to be approximately 16 and 21 MPa m1/2 , respectively. These values far exceed the toughness of monolithic BMG processed by either TS3DP (13 MPa m1/2 ) or casting (5 MPa m1/2 ) [19], and is equivalent to the best nacre-like materials developed previously [25–29] (see Fig. 2f). The fracture toughness of the samples increased as the mass fraction of the SS phase increased, which verified the toughening effect of the SS phase. To exhibit the advance of the TS3DP technique used in this work, we compare the fracture strength and
toughness of the TS3DP-BMGCs with those of other BMGs with similar compositions and those Fe-based BMGCs processed by SLM, SPS and casting (Table S1), as shown in Fig. 2g. Generally, casting and SPS-processed BMGs have high fracture strength but low fracture toughness. The BMGCs with ductile phases (i.e., Cu and CuNi alloy) prepared by SLM have relatively good toughness, but exhibit low strength. Our BMGCs prepared by TS3DP appear at the upper-right regime of the plot, highlighting a good combination of strength and toughness. To essentially uncover the toughening mechanism, the stress fields in the pre-notched BMG composites under three-point bending were visualized in FEM simulation, in comparison with that in monolithic BMG. The contour plots of the third invariant of the stress tensor (determining the volume changes in deformation) in the specimens with equal bending displacement are shown in Fig. 3a-c. The stress level was displayed as a function of the distance from the notch tip (Fig. 3d). A large gradient from tensile stress to compressive stress was observed in the monolithic BMG upon bending. However, as the SS fraction increases, fluctuations in the stress invariant are remarkably reduced within the distance of 350 μm from the notch tip. This would lead to reduction of the driving force of crack propagation. The distribution of Mises stress (Fig. S3) also verified the shielding effect of the SS phase since the stress concentration is effectively mitigated and the stress is widely distributed in the S20 and S50. To clarify the influence of the SS phase on crack propagation, we performed theoretical calculation
116
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117
cal calculation revealed that increase in toughness arose from the unique laminar soft/hard phase configuration which exhibited the effective stress-shielding effect against crack propagation. This approach may provide new options for the design and manufacturing of strong, tough BMGCs for structural applications. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments
Fig. 4. Compressive stress–strain curve of a Fe48 Mo14 Cr15 Y2 C15 B6 BMG composite/FeCoCrNiMn high-entropy alloy (HEA) sample (Ø 2 × 4 mm) prepared via TS3DP technique (the volume fraction of HEA is 80%). Inset is the XRD pattern of the composite, showing a dual structure consisting of both amorphous phase and fcc solid solution.
This work was financially supported by the National Key R&D Program of China (No. 2016YFB1100101), National Natural Science Foundation of China (Grant nos. 51531003, 51771077), the Technical Innovation Project of Hubei Province (2017AAA018) and GDAS’s Project of Science Technology Development (2018GDASCX-0402). The authors are grateful to the Analytical and Testing Center in HUST for technical assistance. Supplementary materials
of the driving force for crack growth in the TS3DP-BMGC according to the inhomogeneous model proposed by Sistaninia and Kolednik [30] (see the details in the Supplementary Materials). A general requirement for crack propagation is that the driving force in front of the crack tip (Jtip ) that exceeds the intrinsic fracture resistance of the material. The calculated Jtip of the crack propagating across the amorphous/SS interface is shown in Fig. 3e. It is seen that Jtip increased sharply as the crack approached interface I, implying that crack would easily penetrate the interface and enter the soft-SS interlayer. Nevertheless, Jtip rapidly decreased to a level far below the far-field J-integral in the amorphous matrix (Jfar ) once the crack entered the SS layer due to the shielding effect. Jtip reached its lowest value at the interface II. In this case, crack tended to deflect and propagate along the intersplat, where the intrinsic resistance to crack propagation was low. To break the next hard-amorphous layer, extra energy is required, thus crack resistance rises steeply. This result indicates that alternative softto-hard materials arrangement plays an essential role in toughening the BMGCs. Similar toughening effect was also achieved in a MG/amorphous-carbon hybrid materials [31]. To present the applicability of the TS3DP technique, we have manufactured several mechanical parts, such as fans and screwdrivers (Fig. 3g, insets). These examples demonstrate that the TS3DP technique is able to fabricate large-sized BMGC component with complex geometries. Moreover, the strategy of preparing nacre-like structure with hard and soft phases using TS3DP can also be adopted to other BMG systems or adding other ductile reinforcements to producing bioinspired hybrid architectures. For instance, we also fabricated a BMG/high-entropy alloy (HEA) dual-structured composite via TS3DP. The fcc FeCoCrNiMn HEA was chosen here owing to its good ductility and damage tolerance [32]. The resultant BMG/HEA composite exhibited a high-strength of > 1.5 GPa and reasonable plasticity of ∼3% (Fig. 4). Hence, this approach may provide new options for design and manufacturing of strong, toughening BMG composites for structural applications. In summary, a strong interfacial bonding with nanoscale compositional gradient between amorphous and SS phases was observed in bioinspired bulk metallic glassy composites prepared by thermal spray additive manufacturing technique. We demonstrated that the soft-SS reinforcement significantly enhanced the mechanical properties of Fe-based BMG. FEM simulations and theoreti-
Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.10. 017. References [1] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [2] A. Inoue, A. Takeuchi, Acta Mater. 59 (2011) 2243–2267. [3] Z. Liu, W. Chen, J. Carstensen, J. Carstensen, J. Ketkaew, R.M.O. Mota, J.K. Guest, J. Schroers, Acta Mater. 105 (2016) 35–43. [4] C. Suryanarayana, A. Inoue, Int. Mater. Rev. 58 (2013) 131–166. [5] M.J. Duarte, J. Klemm, S.O. Klemm, K.J.J. Mayrhofer, M. Stratmann, S. Borodin, A.H. Romero, M. Madinehei, D. Crespo, J. Serrano, S.S.A. Gerstl, P.P. Choi, D. Raabe, F.U. Renner, Science 341 (2013) 372–376. [6] C. Zhang, H. Zhou, L. Liu, Acta Mater. 72 (2014) 239–251. [7] H.X. Li, Z.C. Lu, S.L. Wang, Y. Wu, Z.P. Lu, Prog. Mater. Sci. 103 (2019) 235–318. [8] L.C. Zhang, Z. Jia, F. Lu, S.X. Liang, J. Lu, Prog. Mater. Sci. 105 (2019) 100576. [9] A. Singh, S.R. Paital, A. Andapally, N.B. Dahotre, S.P. Harimkar, Adv. Eng. Mater. 14 (2012) 400–407. [10] J.P. Kelly, S.M. Fuller, K. Seo, E. Novitskaya, V. Eliasson, A.M. Hodge, O.A. Graeve, Mater. Des. 93 (2016) 26–38. [11] S. Pauly, L. Lober, R. Petters, M. Stoica, S. Scudino, U. Kuhn, J. Eckert, Mater. Today 16 (2013) 37–41. [12] X.P. Li, M.P. Roberts, S.O. Keeffe, T.B. Sercombe, Mater. Des. 112 (2016) 217–226. [13] Y.Y. Shen, Y.Q. Li, C. Chen, H.L. Tsai, Mater. Des. 117 (2017) 213–222. [14] M.A. Gibson, N.M. Mykulowycz, J. Shim, R. Fontana, P. Schmitt, A. Roberts, J. Ketkaew, L. Shao, W. Chen, P. Bordeeniothikasem, J.S. Myerberg, R. Fulop, M.D. Verminski, E.M. Sachs, Y.M. Chiang, C.A. Schuh, A.J. Hart, J. Schroers, Mater. Today 21 (2018) 697–702. [15] C. Zhang, X.M. Li, S.Q. Liu, H. Liu, L.J. Yu, L. Liu, J. Alloys Compd. 790 (2019) 963–973. [16] Z. Mahhooba, L. Thorsson, M. Unosson, P. Skoglund, H. West, T. Horn, C. Rock, E. Vogli, O. Harrysson, Appl. Mater. Today 11 (2018) 1–6. [17] C. Yang, C. Zhang, L. Liu, J. Mater. Chem. A 6 (2018) 20992–21002. [18] N. Li, J.J. Zhang, W. Xing, Di. Ouyang, L. Liu, Mater. Des. 143 (2018) 285–296. [19] C. Zhang, W. Wang, Y.C. Li, Y.G. Yang, Y. Wu, L. Liu, J. Mater. Chem. A 6 (2018) 6800–6805. [20] R.P. Wilkerson, B. Gludovatz, J. Watts, A.P. Tomsia, G.E. Hilmas, R.O. Ritchie, Acta Mater. 148 (2017) 147–155. [21] P. Bansal, P.H. Shipway, S.B. Leen, Acta Mater. 55 (2007) 5089–5101. [22] S.J. Wang, T. Jozaghi, I. Karaman, R. Arroyave, Y.I. Chumlyakov, Mater. Sci. Eng. A 694 (2017) 121–131. [23] S.F. Guo, L. Liu, N. Li, Y. Li, Scr. Mater. 62 (2010) 329–332. [24] M.E. Launey, E. Munch, D. Alsem, H.B. Barth, E. Saiz, A.P. Tomsia, R.O. Ritchie, Acta Mater. 57 (2009) 2919–2932. [25] E. Munch, M.E. Launey, D.H. Alsem, E. Saiz, A.P. Tomsia, R.O. Ritchie, Science 322 (2018) 1516–1520. [26] R.P. Wilkerson, B. Gludovatz, J. Watts, A.P. Tomsia, G.E. Hilmas, R.O. Ritchie, Adv. Mater. 28 (2016) 10061–10067. [27] A. Wat, J.I. Lee, C.W. Ryu, B. Gludovatz, J. Kim, A.P. Tomsia, T. Ishikawa, J. Schmitz, A. Meyer, M. Alfreider, D. Kiener, E.S. Park, R.O. Ritchie, Nat. Commun. 10 (2019) 961.
C. Zhang, W. Wang and W. Xing et al. / Scripta Materialia 177 (2020) 112–117 [28] G. Dwivedi, K. Flynn, M. Resnick, S. Sampath, A. Gouldstone, Adv. Mater. 27 (2015) 3073–3078. [29] F. Bouville, E. Maire, S. Meille, V.D. Moortele B, A.J. Stevenson, S. Deville, Nat. Mater. 13 (2014) 508–514. [30] M. Sistaninia, O. Kolednik, Eng. Fract. Mech. 130 (2014) 21–41.
117
[31] W. Wang, C. Zhang, Z.W. Zhang, Y.C. Li, M. Yasir, H.T. Wang, L. Liu, Sci. Rep. 7 (2017) 4084. [32] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, Science 345 (2014) 1153–1158.