Strong and tough Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures

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Here, a set of Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures were fabricated by pressureless infiltration of the Mg melt into ice-templated ceramic scaffolds. The structure and mechanical properties of the composites were elucidated with a special focus on the effects of the types of architectures (lamellar or brick-and-mortar) and matrices (pure Mg or Mg alloy) on the toughening mechanisms. The nacre-like architectures were found to play a role in blunting the cracks via plastic deformation and microcracking, and shielding the cracks from applied stress by promoting crack deflection and uncracked-ligament bridging mechanisms. These composites achieved a good combination of (specific) strength and fracture toughness which are superior to other reported Mg-ceramic and nacre-like metal-ceramic composite materials. Magnesium composites Bioinspired designs Nacre-like architectures Fracture toughness MAX phase Figures Figure 1 Figure 2 Figure 3 Figure 4 1. Introduction Magnesium (Mg) and Mg alloys are promising for weight-critical structural applications owing to their high specific stiffness and specific strength, i.e. , stiffness and strength normalized by density, along with good damping capacity [ 1 ]. However, when compared to other common structural metals, e.g. , steels and aluminum alloys, Mg and Mg alloys are inferior in their stiffness and strength in terms of the absolute values at ambient to elevated temperatures; moreover, they exhibit relatively low fracture toughness at room temperature. A feasible approach to strengthening Mg and Mg alloys is through the introduction of a reinforcement phase into the Mg matrix, in particular by making Mg-ceramic composites [ 2 – 4 ]. In this regard, the MAX phase ceramics appear to be an ideal choice. These ceramics have the general formula of M n+1 AX n (M: early transition metal; A: group A element; X: C and/or N; n = 1–3), are a family of ternary carbides or nitrides with mixed metallic-covalent-ionic atomic bonds. Their characteristics are essentially a combination of the property advantages from both metals, e.g. , good thermal/electrical conductivity and thermal shock resistance, and ceramics, e.g. , high hardness and stiffness, high strength at ambient to elevated temperatures and high corrosion resistance [ 5 – 8 ]; also, they display a good damage tolerance owing to their layered atomic structure [ 9 ]. Moreover, the MAX phase ceramics of Ti-Al-C system have been shown to exhibit good wettability with Mg to form strong interfacial bonding [ 10 – 12 ]. They therefore represent an attractive reinforcement phase for Mg and Mg alloys. There are reports in the literature on the fabrication of such Mg-MAX phase composites to achieve high mechanical strength and good damping properties [ 10 – 14 ]. Nevertheless, despite the lack of data on their fracture toughness, these composites, especially those with relatively high ceramic content exceeding 20 vol.%, were found to fracture in a brittle manner without undergoing visible plastic deformation [ 12 – 14 ]. Indeed, it is a general compromise for metal-ceramic composites that the strength is improved at a cost of fracture toughness [ 15 ]. In the case of Mg-ceramic composites (in bulk form), their fracture toughness normally cannot even reach 10 MPa·m 1/2 when the specific flexural strength exceeds 150 MPa/(g·cm − 3 ); toughnesses above 10 MPa·m 1/2 have only been attained in lower-strength composites [ 16 – 18 ]. Natural nacre, which features a layered structure comprising ~ 95 vol.% aragonite platelets with ~ 5 vol.% organic matter, exhibits much more superior combination of strength and fracture toughness than its constituents [ 19 , 20 ]. Taking the inspiration from nacre offers an effective approach for developing new high-performance composite materials [ 21 – 23 ]. The so-called lamellar and brick-and-mortar architectures are two of the most common structural arrangements mimicking nacre [ 21 , 24 – 29 ]. The stiff and soft constituents are alternately arranged in a layered fashion for the former; whereas the stiff platelets (bricks) are more closely stacked and staggered between each other within the soft matrix (mortar) for the latter. These two architectures also show a large difference in their phase constitution, i.e. , the stiff constituent accounts for a larger volume fraction in the brick-and-mortar architecture than in the lamellar one. Such architectural designs have been implemented in a variety of composite materials, such as those of ceramic-polymer [ 21 , 24 , 25 ], metal-ceramic [ 26 , 27 ] and metal-graphene systems [ 30 , 31 ], and been proven to be effective in toughening materials. However, it is not easy to construct nacre-like architectures, especially the brick-and-mortar architectures, in Mg-based composites. This is largely due to the high reactivity of the Mg melt which tends to react at high temperatures with many of the possible reinforcement phases to form brittle intermetallics, or otherwise display a poor wettability with them [ 3 , 32 ]. Nevertheless, the mechanical properties of nacre-like composites are closely associated with the type of architectures and the characteristics of matrices [ 21 , 24 – 26 ], but these factors and their effects remain to be explored for Mg-based composites. 2. Results Figure 1 shows the microstructures and phase constitution of the infiltrated Mg-MAX phase composites. The constituents that were respectively rich in metal and ceramic phases showed an alternate arrangement in a layer fashion in the composites (Fig. 1 a,b), which is qualitatively similar to the structure of natural nacre. The ceramic-rich constituents were interconnected between layers, via bridges for the lamellar architecture and by overlapping for the brick-and-mortar one. As compared to the lamellar architecture, the ceramic-rich constituents in the brick-and-mortar architecture were more closely stacked and accounted for a larger volume fraction in the composite. In both architectures, the ceramic-rich constituents were indeed composed of both metal and ceramic phases instead of monolithic ceramics (Fig. 1 c). The metal phase within these constituents had ultrafine grains typically smaller than 2 µm (Fig. 1 d), due to the constraint of ceramic phase on the growth of metal grains during the solidification process after the melt infiltration procedure. The ceramic contents in the composites with pure Mg and AZ91D alloy as the metal component were calculated to be ~ 11.4 vol.% and ~ 14.0 vol.% for the lamellar architecture, and ~ 32.6 vol.% and ~ 32.5 vol.% for the brick-and-mortar architecture, respectively, according to the rule-of-mixtures based on their measured densities (respectively, ~ 2.02 g·cm − 3 and ~ 2.16 g·cm − 3 for the lamellar architecture, and ~ 2.54 g·cm − 3 and ~ 2.60 g·cm − 3 for the brick-and-mortar architecture). The volume fractions of the ceramic-rich constituents were determined by image analysis to be ~ 31.1% and ~ 51.3% for the pure Mg composites with lamellar and brick-and-mortar architectures, respectively; the corresponding values for the AZ91D alloy composites were ~ 30.6% and ~ 52.0%. As such, the volume fractions of the metal phase within the ceramic-rich constituents were derived to be roughly 63.5% and 54.2% for the pure Mg and AZ91D alloy composites with the lamellar architecture, and to be 36.5% and 37.4% for those with the brick-and-mortar architecture, respectively. EDS analysis revealed that the Mg element and the elemental Ti, Al and C were principally concentrated in the metal and ceramic-rich constituents (Fig. 1 e and Figure S1 in Supplementary Materials), respectively. This, combined with the XRD results (Fig. 1 f), clearly indicates that the composites largely retained the pre-designed phase constitution of Mg and Ti 3 AlC 2 phases after melt infiltration. In addition, a small amount of MgO and Mg 17 Al 12 phases were also recognized in the XRD patterns. The former should be the product of the reaction between the Mg melt and the surface oxides of ceramic phase during the melt infiltration process (the oxygen content in the ball-milled Ti 3 AlC 2 platelets was measured to be ~ 4.2 wt.%). The Mg 17 Al 12 phase may have formed due to the interfacial reaction between the melting Mg and Ti 3 AlC 2 phase [ 12 , 14 ]. Similar reactions have been identified between other metals and MAX phases [ 33 , 34 ], e.g. , in Ag-Ti 3 AlC 2 system [ 33 ], and been demonstrated to play a role in promoting a strong interfacial bonding between the metal and ceramic phases. The nacre-like Mg-MAX phase composites exhibit large differences in their local mechanical properties. As shown in Fig. 2 a, for the lamellar architecture subject to an equal load under nanoindentation condition, relatively larger displacements were obtained: (i) in the metal constituent (inter-lamellar region) than in the ceramic-rich constituent (lamella) for the same type of metal component, or (ii) in the composites with pure Mg than AZ91D alloy than in the metal component for the same regions. Specifically, the nanoindentation hardness of the ceramic-rich constituents was markedly higher than that of the inter-lamellar region respectively by ~ 1.7 and ~ 2.1 times in the pure Mg and AZ91D alloy composites (Fig. 2 b). Additionally, the lamella and inter-lamellar regions of the AZ91D alloy composites were harder by ~ 1.4 and ~ 1.1 times, respectively, than in the pure Mg composites. However, the local elastic moduli determined by nanoindentation testing were similar between the composites containing the two types of metal component, both measured to be ~ 91.4 GPa and ~ 61.9 GPa, respectively, for the lamella and inter-lamellar regions. The global mechanical properties of the composites are also related to the types of nacre-like architecture and metal component. For the same type of metal component, the brick-and-mortar architecture exhibited higher flexural strengths than the lamellar one, respectively up to 491.1 ± 27.4 MPa and 652.4 ± 45.9 MPa for the pure Mg and AZ91D alloy composites (Fig. 2 c). Nevertheless, this was accompanied by an obvious compromise in the ductility for both metal components. With respect to the effects of metal component, the AZ91D alloy composites invariably displayed relatively higher flexural strengths when considering the same type of architecture, whereas the pure Mg composites displayed better ductility. All these composites showed microscopically staircase-like morphologies on their fracture surfaces, which are consistent with their layered structure (Fig. 2 d-f). Apparent dimples that are indicative of plastic deformation were formed in the metal constituent for the pure Mg composites with lamellar architecture (Fig. 2 d,e). Such feature became much less evident in the AZ91D alloy composites (Fig. 2 f), consistent with their reduced ductility. Figure 3 a shows the representative single-edge notched bending load-displacement curves of the nacre-like Mg-MAX phase composites. Instantaneous fracture occurred at the maximum bending load without stable crack propagation in the AZ91D alloy composites with the brick-and-mortar architecture. In comparison, the bending load decreased with increasing displacement after reaching its maximum in the other composites. In particular, the pure Mg composites with the lamellar architecture showed gradual load drop along with a large displacement, indicating a stable cracking process. The fracture resistance of the composites in terms of \(J\) -integral and equivalent stress intensity \({K}_{J}\) as a function of crack extension ( \(\varDelta a\) ) are shown in Fig. 3 b,c. The AZ91D alloy composites with brick-and-mortar architecture fractured abruptly at the onset of crack propagation, giving \(J\) and \(K\) of 2.1 ± 0.7 kJ·m − 2 and 14.4 ± 2.7 MPa·m 1/2 , respectively. However, all the other composites, i.e. , the AZ91D alloy composites with lamellar architecture and the pure Mg composites, displayed a stable increase in \(J\) and \({K}_{J}\) with crack extension, showing rising crack resistance-curve (R-curve) behavior. Higher fracture toughness was generated in the lamellar architecture for both types of metal component, or in the pure Mg composites for a given type of architecture. In particular, the pure Mg composites with lamellar architecture exhibited the highest fracture toughness, with the critical \(J\) and \({K}_{J}\) up to respectively 15.9 ± 2.5 kJ·m − 2 and 32.5 ± 2.6 MPa·m 1/2 , which were determined by the maximum crack extension of ~ 0.5 mm according to the ASTM Standard E1820 (as indicated by the dashed lines) [ 35 ]. The nacre-like architectures have been proven to be effective in activating a series of extrinsic toughening mechanisms for toughening materials, such as crack deflection, uncracked-ligament bridging and frictional sliding between crack faces [ 15 , 21 ]. These mechanisms were clearly introduced into the Mg-MAX phase composites by the lamellar architecture, as manifested by the obviously tortuous cracking paths (Fig. 3 d); nevertheless, such behavior became much less evident for the brick-and-mortar architecture. With regard to the effects of the type of metal component, obvious plastic deformation occurred in the inter-lamellar region near the primary crack in the pure Mg composites with lamellar architecture; whereas microcracks were formed in the ceramic-rich constituents (Fig. 3 e). In comparison, for the AZ91D alloy composites, both lamella and inter-lamellar regions exhibited apparent microcracking without visible plastic deformation (Fig. 3 f). 3. Discussion Both plastic deformation and microcracking in the vicinity of the crack tip, which were dominant respectively in pure Mg and AZ91D alloy composites, offer a means for dissipating mechanical energy and can promote crack bifurcation (Figure S2). These mechanisms, which can be termed intrinsic toughening mechanisms [ 15 , 36 ], are principally active ahead of the crack tip to blunt the crack tip. In addition, they play a role in promoting the crack deflection following the slip bands in inter-lamellar regions or along the microcracks (Figure S2) and inducing the formation of uncracked-ligament bridges, i.e. , between the main cracks and voids or microcracks. These mechanisms, which can be termed extrinsic toughening mechanisms [ 15 , 36 ], are principally active at or behind the crack tip to reduce the crack-driving force actually experienced at the crack tip, thereby shielding the crack tip from applied stress. The effectiveness of the above toughening mechanisms in the nacre-like Mg-MAX phase composites is associated with the types of architectures (lamellar or brick-and-mortar) and metal components (pure Mg or AZ91D alloy). The lamellar architecture, which contains a larger volume fraction of metal constituent than the brick-and-mortar one, enables a more stable crack extension via crack defection, bifurcation and uncracked-ligament bridging mechanisms, leading to higher crack-growth toughness after the onset of cracking. Such trend is opposite to that found in nacre-like composites of ceramic-polymer (polymethylmethacrylate or epoxy resin) [ 21 , 25 ] and alumina-metallic glass [ 26 ] systems where the brick-and-mortar architecture is tougher than the lamellar one. Indeed, the relatively soft constituents (of polymers or metallic glass) that comprise the “mortar” appear to be brittle and can hardly undergo large plastic deformation, which is different from the case for pure Mg and AZ91D alloy. As such, the fracture resistance in these composites is mainly developed from the extrinsic toughening mechanisms which are more effective in the brick-and-mortar architecture owing to its higher density of interfaces and bridges. With respect to the effects of the metal component, pure Mg exhibits a lower strength than the AZ91D alloy in the as-cast state (Figure S3), and thereby can be more easily deformed in the Mg-MAX phase composites. Such easy inelastic deformation of the soft phase, as compared to the fracture of stiff constituents, is essential for ensuring the toughening effects of the nacre-like architectures in promoting both the intrinsic and extrinsic mechanisms [ 15 , 21 ]. Figure 4 a presents a comparison of the fracture toughness and flexural strength for the Mg-MAX phase composites with their component materials of pure Mg, AZ91 (Mg-Al-Zn) series alloys and bulk Ti 3 AlC 2 ceramics [ 5 , 6 , 16 , 17 , 37 – 40 ]. Most MAX phase ceramics were reported to exhibit a flexural strength of less than 560 MPa (although they can be strengthened by fine modulation of their microstructures [ 41 ]); and their fracture toughness rarely exceeds 10 MPa·m 1/2 [ 5 , 6 ]. Pure Mg and AZ91 series alloys generally exhibit a fracture toughness of lower than 24 MPa·m 1/2 and a flexural strength below 400 MPa [ 37 – 40 ]. In comparison, the flexural strength of the nacre-like Mg-MAX phase composites is comparable to, or even higher than, those of monolithic Ti 3 AlC 2 ceramics, while their fracture toughness is comparable to that of the metal phase. This indicates the achievement of a good combination of property advantages in the composites inherited from their constituent materials. Figure 4 b shows the variation in the flexural strength as a function of the ceramic content in Mg-based composites reinforced by different types of ceramics [ 16 , 17 , 38 , 39 , 42 – 44 ]. The current Mg-MAX phase composites exhibit relatively higher flexural strength than most other materials at equal ceramic contents. Figure 4 c shows a comparison of the specific fracture toughness and specific flexural strength, i.e. , fracture toughness and flexural strength normalized by density, for the nacre-like composites of different metal-ceramic systems fabricated by ice-templating and subsequent melt infiltration procedures [ 17 , 26 , 45 – 49 ]. The Mg-MAX phase composites show a high specific flexural strength ranging from ~ 200 MPa/(g·cm − 3 ) to over 250 MPa/(g·cm − 3 ) combined with a high specific fracture toughness. In particular, their specific fracture toughness exceeds those of most other nacre-like metal-ceramic composites with the same level of specific flexural strength. This good combination of properties makes the nacre-like Mg-MAX phase composites appealing for many lightweight structural applications. In summary, a set of Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures were fabricated by pressureless infiltration of pure Mg and AZ91D alloy melt into ice-templated porous Ti 3 AlC 2 scaffolds. The ceramic-rich constituents were alternately arranged with the metals in a layered fashion in the composites, and were more densely stacked in the brick-and-mortar architecture than in the lamellar one. They had an ultrafine structure comprising both ceramic and metal phases that were interconnected between adjacent layers. The mechanical properties of the Mg-MAX phase composites were associated with the types of architectures and metal components. Specifically, higher flexural strength, yet accompanied with inferior ductility, was generated in the AZ91D (Mg-Al-Zn-Mn) magnesium alloy composites for a given architecture or in the brick-and-mortar architecture for a given metal component. Apart from the AZ91D alloy composites with brick-and-mortar architecture, the Mg-MAX phase composites displayed rising R-curve behavior with stable crack propagation process. Crack advance was seen to involve bifurcation with crack blunting via plastic deformation and microcracking, which were respectively dominant in pure Mg and AZ91D alloy composites; additionally, these materials were extrinsically toughened by crack deflection and uncracked-ligament bridging resulting in significant crack-tip shielding, although these latter mechanisms were more evident in the lamellar than in the brick-and-mortar architectures. The composites achieved an effective combination of the property advantages from their component materials. They were stronger than most other Mg-ceramic composites at equal ceramic contents and exhibited a good combination of specific fracture toughness and specific flexural strength among nacre-like metal-ceramic composites, thereby demonstrating a potential for lightweight structural applications. We believe that his study may further give insights for the design of new high-performance Mg-based composites. 4. Methods Commercial Ti 3 AlC 2 powders (400 mesh, Haixin metal, China) were ball milled in an ethanol dispersant in a mixed gas of 20 vol.% oxygen and 80 vol.% argon at a speed of 500 rpm for 48 h. In this process, the MAX phase powders were delaminated into ultrafine platelets along their relatively weak MX/A interlayer bonding of layered structure [ 50 ]. The equivalent particle size of the platelets was measured to be 100–300 nm using laser diffraction granulometry. The platelets were dried in air for 2 days and then were dispersed in deionized water at a mass ratio of 1:2 under ultrasonic treatment at 50 W for 0.5 h. The slurry was mixed with an anionic dispersant (Darvan CN, R.T. Vanderbilt, USA), polyvinyl alcohol (Meryer, China) and hydroxypropyl methylcellulose (Meryer, China), which accounted, respectively, for 0.2 wt.%, 1.5 wt.% and 0.5 wt.% of the platelets. These additives were used for promoting dispersion, increasing the viscosity of slurry to mitigate gravitational sedimentation, and binding the platelets after the removal of the water. The slurry was ball milled at a speed of 25 rpm for 24 h, and then was poured into a cuboid Teflon mold with an inner square cross-section of 30 mm × 30 mm which was fixed onto a customized freezing apparatus. The mold was designed to have non-uniform lateral walls which was thinner at one side (3 mm) than the other three sides (20 mm), and was sealed using a wedge with a slope angle of 25° at its bottom with the thinner end of the wedge conforming to the thinner wall of the mold. Such designs were implemented to create a horizontal temperature gradient in addition to the vertical one in the slurry during the freezing process [ 51 , 52 ]. The mold was descended at a constant rate of 0.6 mm·min − 1 to immerse into an insulated box containing liquid nitrogen until the slurry was fully frozen. The frozen body was demolded and freeze-dried in a vacuum below 5 Pa for at least 72 h using a Scientz-10ND freeze drier (Scientz Biotechnology, China). Porous Ti 3 AlC 2 scaffolds with lamellar architecture were obtained by sintering the freeze-dried samples in flowing argon gas at 900°C for 1 h. Some of the scaffolds were infiltrated with paraffin wax and then uniaxially pressed along the normal direction of lamellae at 75°C (~ 15°C higher than the melting point of paraffin wax) under a pressure of 20 ~ 35 MPa for 2 h for densification. These scaffolds were re-sintered in flowing argon gas at 950°C for 1 h to create interconnections between adjacent lamellae or bricks. Pure Mg with a purity of 99.99 wt.% and a widely used AZ91D cast Mg alloy (with ~ 8.5–9.5 wt.% Al, ~ 0.45–0.9 wt.% Zn and ~ 0.17–0.5 wt.% Mn) were used for fabricating the Mg-MAX phase composites. The sintered Ti 3 AlC 2 scaffolds were infiltrated with the melt of pure Mg or AZ91D alloy in flowing argon gas at 850°C, some 250°C higher than their melting points, for 1.5 h. The densities of the infiltrated composites were measured using the Archimedes’ method [ 53 ]. The phase constitution was determined by X-ray diffraction (XRD) using a Bruker D8 Advance X-ray diffractometer (Bruker AXS, Germany) with Cu-Kα radiation. Scanning electron microscopy (SEM) imaging was conducted using an Inspect F50 field-emission scanning electron microscope (FEI, USA) operating at an accelerating voltage of 20 kV. Energy dispersive X-ray spectroscopy (EDS) measurements were carried out using a Bruker EDS system attached to the microscope. Scanning transmission electron microscopy (STEM) imaging was performed using a Tacnai-G2 F30 transmission electron microscope (FEI, USA) operating at an accelerating voltage of 300 kV. Nanoindentation tests were performed using a G200 nano indenter (KLA, USA) with a Berkovich diamond tip at a constant loading rate of 0.67 mN·s − 1 to a peak load of 10 mN with a holding time of 10 s followed by unloading. Three-point bending tests were conducted using an Instron E1000 testing system (Instron, USA) at room temperature with a constant displacement rate of 0.1 mm·min − 1 on beam samples with dimensions of 2 mm × 1.5 mm × 25 mm and a loading span of 20 mm, in line with the ASTM Standard C1161 [ 54 ]. Single-edge notched bending tests were performed on a JEOL MicroTest stage inside the chamber of a JEOL JSM-6510 scanning electron microscope (JEOL, Japan) at a displacement rate of 5 µm·min − 1 . Samples with a width of 4 mm, thickness of 2 mm and loading span of 16 mm were side notched to ~ 2 mm in depth using a low-speed diamond wire saw with the notch sharpened to a tip radius of 5 ~ 10 µm using a razor blade with 1 µm diamond paste, in general accordance with the ASTM Standard E1820 [ 35 ]. The loads in both three-point and singe-edge notched bending tests were perpendicular to the layered structure of the composites, which is consistent with the general loading configuration for natural nacre. The fracture morphologies of samples were characterized by SEM imaging. Nonlinear-elastic fracture mechanics methods were used to evaluate the fracture toughness of the composites in terms of their J -integral. The J values were calculated from the applied load and instantaneous crack length by considering the elastic ( \({J}_{el}\) ) and plastic ( \({J}_{pl}\) ) components, i.e. , \(J={J}_{el}+{J}_{pl}\) , according to the ASTM Standard E1820 [ 35 ]. \({J}_{el}\) was formulated using the linear-elastic stress intensity factor K according to \({J}_{el}=\frac{{K}^{2}\left(1-{v}^{2}\right)}{E}\) , where \(E\) is the Young’s modulus and \(v\) is the Poisson’s ratio. The \(E\) values were approximated using the averages of those calculated following the rule-of-mixtures under the iso-strain (Voigt model) and iso-stress (Reuss model) conditions [ 55 ]. In view that the mode I stresses were parallel to the layered structure of the nacre-like composites, such approach was expected to give a conservative approximation for the Young’s modulus and therefore the fracture toughness. The Poisson’s ratio \(v\) of the composites was approximated in the same fashion to be 0.28 and 0.26, respectively, for the lamellar and brick-and-mortar architectures. The plastic component \({J}_{pl}\) was calculated from the plastic area ( \({A}_{pl}\) ) under the load-displacement curve according to \({J}_{pl}=\frac{1.9{A}_{pl}}{Bb}\) [ 35 ], where \(B\) is the thickness of sample and \(b\) is the length of the uncracked ligament. The equivalent \(K\) -based stress intensity was obtained from the J -integral based on the standard \(J\) - \(K\) equivalence as \({K}_{J}={\left(\frac{JE}{\left(1-{v}^{2}\right)}\right)}^{1/2}\) in the case of mode I fracture [ 35 ]. Declarations Data availability The data that support the findings of this study are available from the corresponding author, Prof. Zengqian Liu, at [email protected] , upon reasonable request. Acknowledgements The authors are grateful for the financial support by the National Key R&D Program of China (2020YFA0710404), the National Natural Science Foundation of China (52173269 and 51871216), the KC Wong Education Foundation (GJTD-2020-09), the Youth Innovation Promotion Association CAS, and the CAS Project for Young Scientists in Basic Research (YSBR-025). ROR was supported by the Multi-University Research Initiative (AFOSR-FA9550-15-1-0009) from the Air Force Office of Scientific Research. 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Mineral nano-interconnectivity stiffens and toughens nacre-like composite materials. Adv Mater 2017;29:1605039. Tan GQ, Zhang J, Zheng L, Jiao D, Liu ZQ, Zhang ZF, Ritchie RO. Natural-inspired nacre-like composites combing human tooth-matching elasticity and hardness with exceptional damage tolerance. Adv Mater 2019;31:1904603. Wat A, Lee JI, Ryu CW, Gludovatz B, Kim JY, Tomsia AP, Ishikawa T, Schmitz J, Meyer A, Alfreider M, Kiener D, Park ES, Ritchie RO. Bioinspired nacre-like alumina with a bulk-metallic glass-forming alloy as a compliant phase. Nat Commun 2019;10:961. Ferraro C, Meille S, Rethore J, Ni N, Chevalier J, Saiz E. Strong and tough metal/ceramic micro-laminates. Acta Mater 2018;144:202-215. Zhang MY, Zhao N, Yu Q, Liu ZQ, Qu RT, Zhang J, Li SJ, Ren DC, Berto F, Zhang ZF, Ritchie RO. On the damage tolerance of 3-D printed Mg-Ti interpenetrating-phase composites with bioinspired architectures. Nat Commun 2022;13:3247. Morsali S, Qian D, Minary-Jolandan M. Designing bioinspired brick-and-mortar composites using machine learning and statistical learning. Commun Mater 2022;1(1):1-11. Xiong DB, Cao M, Guo Q, Tan ZQ, Fan GL, Li ZQ, Zhang D. Graphene-and-copper artificial nacre fabricated by a preform impregnation process: bioinspired strategy for strengthening-toughening of metal matrix composite. ACS Nano 2015;9:6934-6943. Zhang X, Zhao NQ, He CN. The superior mechanical and physical properties of nanocarbon reinforced bulk composites achieved by architecture design – a review. Prog Mater Sci 2020;113:100672. Yao YT, Chen LQ. Processing of B 4 C particulate-reinforced magnesium-matrix composites by metal-assisted melt infiltration technique. J Mater Sci Technol 2014;30:661-665. Ding KK, Zhang KG, Ding JX, Zhang X, Xia XX, Chen LM, Ran SL, Ma CJ, Yang L, Zhang PG, Sun ZM. Effect of Al atomic layer on the wetting behavior, interface structure and electrical contact properties of silver reinforced by Ti 3 AlC 2 ceramic. Ceram Int 2022;48:190-198. Wang DD, Tian WB, Lu CJ, Ding JX, Zhu YF, Zhang M, Zhang PG, Sun ZM. Comparison of the interfacial reactions and properties between Ag/Ti 3 AlC 2 and Ag/Ti 3 SiC 2 electrical contact materials. J Alloys Compd 2021;857:157588. ASTM Standard E1820-13, Standard Test Method for Measurement of Fracture Toughness , American Society for Testing and Materials International, West Conshohocken, PA, USA 2013. Kruzic JJ, Nalla RK, Kinney JH, Ritchie RO. Crack blunting, crack bridging and resistance-curve fracture mechanics in dentin: effect of hydration. Biomaterials 2003;24:5209-5221. Somekawa H, Mukai T. Effect of grain refinement on fracture toughness in extruded pure magnesium. Scr Mater 2005;53:1059-1064. Pitchayyapillai G, Mohamed MJS, Dhanraj G, Prince RMR, Rajeshwaran M, Mangrulkar A. Influence of B 4 C on mechanical properties of AZ91 magnesium matrix composites. Mater Today: Proc 2022;59:1438-1441. Park IM, Choi LD, Cho KM. Mechanical and fracture behaviors of (Al 2 O 3 +SiC p )/AZ91 hybrid Mg matrix composite. Mater Sci Forum 2004;449:653-656. Ashby M. Mapping the fracture properties of engineering materials. Philos Mag 2013;93:3878-3892. Wang WJ, Li CW, Zhai HX, Wang CG. Preparation of high-strength Ti 3 AlC 2 by spark plasma sintering. Appl Ceram Technol 2015;12:E126-E131. Sohn KS, Euh K, Lee S, Park I. Mechanical property and fracture behavior of squeeze-cast Mg matrix composites. Metall Mater Trans A 1998;29:2543-2554. Rangaraj L, Sagar RV, Stalin M, Raghavendra K, Venkateswarlu K. Synthesis, characterization, and mechanical properties evaluation of Mg-Ti 3 AlC 2 composites produced by powder metallurgy/hot pressing. Metall Mater Trans A 2019;50:3714-3723. Chen J, Bao CG, Chen WH, Zhang L, Liu JL. Mechanical properties and fracture behavior of Mg-Al/AlN composites with different particle contents. J Mater Sci Technol 2017;33:668-674. Li YL, Guo RF, Hu ZJ, Shen P. Construction of nacre-mimetic composites with a ‘‘brick-and-mortar” architecture based on structural defects in ice-templating. Mater Des 2021;204:109668. Launey ME, Munch E, Alsem GH, Saiz E, Tomsia AP, Ritchie RO. A novel biomimetic approach to the design of high-performance ceramic-metal composites. J R Soc Interface 2010;7:741-753. Liu Q, Ye F, Gao Y, Liu SC, Yang HX, Zhou ZQ. Fabrication of a new SiC/2024Al co-continuous composite with lamellar microstructure and high mechanical properties. J Alloys Compd 2014;585:146-153. Sun MQ, Shen P, Jiang QC. Microstructures and mechanical characterizations of high-performance nacre-inspired Al/Al 2 O 3 composites. Compos Part A Appl Sci Manuf 2019;121:465-473. Wang Y, Liu Q, Zhang B, Zhang HQ, Jin YC, Zhong ZX, Ye J, Ren YH, Ye F, Wang W. High damage-tolerance bio-inspired B 4 C/2024Al composites with adjustable mechanical performance by tuning ceramic thickness. Mater Sci Eng A 2021;819:141469. Li XQ, Xie X, Julian JG, Malzbender J, Yang R. Mechanical and oxidation behavior of textured Ti 2 AlC and Ti 3 AlC 2 MAX phase materials. J Eur Ceram Soc 2020;40:5258-5271. Bai H, Walsh F, Gludovatz B, Delattre B, Huang CL, Chen Y, Tomsia AP, Ritchie RO. Bioinspired hydroxyapatite/poly (methyl methacrylate) composite with a nacre-mimetic architecture by a bidirectional freezing method. Adv Mater 2015;28:50-56. Qiu WX, Zhang J, Tan GQ, Gao KF, Zhang MY, Liu ZQ, Zhang ZF. Continuous ice-templating of macro-porous materials with uniformly ordered architecture. Sci China Mater. doi:10.1007/s40843-022-2078-8. Hughen SW. Archimedes revisited: a faster, better, cheaper method of accurately measuring the volume of small objects. Phys Educ 2005;40:468-474. ASTM Standard C1161-18, Standard Test Method for Flexural Strength of Advanced Ceramics at Ambient Temperature , American Society for Testing and Materials International, West Conshohocken, PA, USA 2018. Watt JP, Davies GF, O’Connell RJ. The elastic properties of composite materials. Rev Geophys Space Phys 1976;14:541-563. Additional Declarations There is NO Competing Interest. Supplementary Files SupplementaryMaterialsror.docx Cite Share Download PDF Status: Published Journal Publication published 09 May, 2023 Read the published version in Communications Materials → Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-1896279","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":124256783,"identity":"af6a238e-0e89-4d1b-9c89-6f2fc0cebb27","order_by":0,"name":"Yanyan Liu","email":"","orcid":"","institution":"Institute of Metal Research, Chinese Academy of Sciences","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Yanyan","middleName":"","lastName":"Liu","suffix":""},{"id":124256784,"identity":"0352e2c0-fa53-48ba-a2db-d07efb064e68","order_by":1,"name":"Xi Xie","email":"","orcid":"","institution":"Institute of Metal Research, Chinese Academy of Sciences","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Xi","middleName":"","lastName":"Xie","suffix":""},{"id":124256785,"identity":"4b2b072e-71c3-4c82-b996-8ff860c0282f","order_by":2,"name":"Zengqian Liu","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA4ElEQVRIiWNgGAWjYNACAwYGPmYGxgcJJGlhY2ZgNiBBCxCwAZEEUSrl23sPv3hTYJfHxs5jVvEw5x4Df/sBxs8FeLQw9pxLs5xjkFzMxsxjdiNxWzGDxJkEZukZeLQwS+SYGfMYMCe2QbQkMDDcAPqLB58X5N+AtNSDtRSAtMgT0sIjwWP8mMfgMFgLA0iLASEtEjw5ZoxzDI4DtbAVSwC18BieSWyWxqdFvv2M8Yc3f6oT+/kPb/z4c1uCnNzxwwc/49MC8o4EsgIgm7EBvwZgoH0gYOYoGAWjYBSMdAAArvQ+PzD8sekAAAAASUVORK5CYII=","orcid":"","institution":"Institute of Metal Research, Chinese Academy of Sciences","correspondingAuthor":true,"submittingAuthor":false,"prefix":"","firstName":"Zengqian","middleName":"","lastName":"Liu","suffix":""},{"id":124256786,"identity":"f6760fd9-3548-4871-9307-94fbcf442836","order_by":3,"name":"Qin Yu","email":"","orcid":"https://orcid.org/0000-0002-5099-0327","institution":"","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Qin","middleName":"","lastName":"Yu","suffix":""},{"id":124256787,"identity":"92445a5f-235d-4ed1-9703-f28dc0cd2586","order_by":4,"name":"Xuegang Wang","email":"","orcid":"","institution":"","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Xuegang","middleName":"","lastName":"Wang","suffix":""},{"id":124256788,"identity":"ae364785-5c62-4ce0-b360-350583161cac","order_by":5,"name":"Qing Jia","email":"","orcid":"","institution":"","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Qing","middleName":"","lastName":"Jia","suffix":""},{"id":124256789,"identity":"5780383b-2f91-401b-9c7e-67e3fb9a5dc2","order_by":6,"name":"Zhefeng Zhang","email":"","orcid":"","institution":"","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Zhefeng","middleName":"","lastName":"Zhang","suffix":""},{"id":124256790,"identity":"f7a0f7df-1db7-4b4a-9ce6-5628627cac6c","order_by":7,"name":"Rui Yang","email":"","orcid":"","institution":"","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Rui","middleName":"","lastName":"Yang","suffix":""},{"id":124256791,"identity":"a4f9b484-dbd3-4ab2-b876-84371b82cd4a","order_by":8,"name":"Robert Ritchie","email":"","orcid":"https://orcid.org/0000-0002-0501-6998","institution":"University of California Berkeley","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Robert","middleName":"","lastName":"Ritchie","suffix":""}],"badges":[],"createdAt":"2022-07-26 06:21:02","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-1896279/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-1896279/v1","draftVersion":[],"editorialEvents":[{"content":"https://doi.org/10.1038/s43246-023-00358-3","type":"published","date":"2023-05-09T04:00:00+00:00"}],"editorialNote":"","failedWorkflow":false,"files":[{"id":24528989,"identity":"ad1ca815-33a8-474b-a007-b5773c589a40","added_by":"auto","created_at":"2022-07-29 18:55:44","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":1202929,"visible":true,"origin":"","legend":"\u003cp\u003eMicrostructures and phase constitution of the nacre-like Mg-MAX phase composites. (a, b) SEM micrographs of the Mg-MAX phase composites with nacre-like lamellar (a) and brick-and-mortar (b) architectures. (c, d) Magnified SEM (c) and STEM (d) micrographs of the ceramic-rich constituent in the composites showing its ultrafine composite structure (c: pure Mg composite with lamellar architecture; d: AZ91D alloy composite with brick-and-mortar architecture). (e) Distribution of elements measured by EDS in the AZ91D alloy composite with lamellar architecture. (f) XRD patterns of the Mg-MAX phase composites containing pure Mg and AZ91D alloy as the metal component (by taking the brick-and-mortar architecture as an example).\u003c/p\u003e","description":"","filename":"floatimage1.png","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/c0793cb4710c1d83b078cff6.png"},{"id":24529130,"identity":"2a91d8db-3eff-40ca-a110-617ec3c0621d","added_by":"auto","created_at":"2022-07-29 19:00:44","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1144738,"visible":true,"origin":"","legend":"\u003cp\u003eMechanical properties of the nacre-like Mg-MAX phase composites. (a, b) Representative load-displacement curves (a) and the hardness and elastic modulus (b) measured by nanoindentation testing for the metal (inter-lamellar) and ceramic-rich (lamella) constituents in the pure Mg and AZ91D alloy composites with lamellar architecture. (c) Flexural stress-strain curves of the Mg-MAX phase composites with different architectures and metal components. (d-f) SEM micrographs of the fracture surfaces for the pure Mg (d, e) and AZ91D alloy (f) composites with lamellar architecture after bending. (e) is a magnified view of the dashed box in (d).\u003c/p\u003e","description":"","filename":"floatimage2.png","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/bb63865e8728f074dfc5535c.png"},{"id":24528990,"identity":"9f67f93d-87bc-4eeb-92c9-8386bc1a0c40","added_by":"auto","created_at":"2022-07-29 18:55:44","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":265074,"visible":true,"origin":"","legend":"\u003cp\u003eSee image above for figure legend\u003c/p\u003e","description":"","filename":"floatimage3.png","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/d2bd20d0498b617b3763a025.png"},{"id":24528991,"identity":"dd865a08-8bff-442b-95a7-ff7effa12e50","added_by":"auto","created_at":"2022-07-29 18:55:44","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":220218,"visible":true,"origin":"","legend":"\u003cp\u003eSee image above for figure legend\u003c/p\u003e","description":"","filename":"floatimage4.png","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/32bddcfa7286d56932d75d8a.png"},{"id":36907522,"identity":"5f69fc3b-76fe-40b1-8a7e-4b948ca03588","added_by":"auto","created_at":"2023-05-12 02:46:19","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2185981,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/e19c8c93-af50-43ac-a7c1-c67d6e4a01b0.pdf"},{"id":24528993,"identity":"e84bdcdd-7c46-4881-bccf-2230d2255a41","added_by":"auto","created_at":"2022-07-29 18:55:44","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":581331,"visible":true,"origin":"","legend":"","description":"","filename":"SupplementaryMaterialsror.docx","url":"https://assets-eu.researchsquare.com/files/rs-1896279/v1/2a499533e9d12640e7019568.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Strong and tough Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eMagnesium (Mg) and Mg alloys are promising for weight-critical structural applications owing to their high specific stiffness and specific strength, \u003cem\u003ei.e.\u003c/em\u003e, stiffness and strength normalized by density, along with good damping capacity [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e]. However, when compared to other common structural metals, \u003cem\u003ee.g.\u003c/em\u003e, steels and aluminum alloys, Mg and Mg alloys are inferior in their stiffness and strength in terms of the absolute values at ambient to elevated temperatures; moreover, they exhibit relatively low fracture toughness at room temperature. A feasible approach to strengthening Mg and Mg alloys is through the introduction of a reinforcement phase into the Mg matrix, in particular by making Mg-ceramic composites [\u003cspan additionalcitationids=\"CR3\" citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. In this regard, the MAX phase ceramics appear to be an ideal choice. These ceramics have the general formula of M\u003csub\u003en+1\u003c/sub\u003eAX\u003csub\u003en\u003c/sub\u003e (M: early transition metal; A: group A element; X: C and/or N; n\u0026thinsp;=\u0026thinsp;1\u0026ndash;3), are a family of ternary carbides or nitrides with mixed metallic-covalent-ionic atomic bonds. Their characteristics are essentially a combination of the property advantages from both metals, \u003cem\u003ee.g.\u003c/em\u003e, good thermal/electrical conductivity and thermal shock resistance, and ceramics, \u003cem\u003ee.g.\u003c/em\u003e, high hardness and stiffness, high strength at ambient to elevated temperatures and high corrosion resistance [\u003cspan additionalcitationids=\"CR6 CR7\" citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]; also, they display a good damage tolerance owing to their layered atomic structure [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e]. Moreover, the MAX phase ceramics of Ti-Al-C system have been shown to exhibit good wettability with Mg to form strong interfacial bonding [\u003cspan additionalcitationids=\"CR11\" citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e]. They therefore represent an attractive reinforcement phase for Mg and Mg alloys.\u003c/p\u003e \u003cp\u003eThere are reports in the literature on the fabrication of such Mg-MAX phase composites to achieve high mechanical strength and good damping properties [\u003cspan additionalcitationids=\"CR11 CR12 CR13\" citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. Nevertheless, despite the lack of data on their fracture toughness, these composites, especially those with relatively high ceramic content exceeding 20 vol.%, were found to fracture in a brittle manner without undergoing visible plastic deformation [\u003cspan additionalcitationids=\"CR13\" citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. Indeed, it is a general compromise for metal-ceramic composites that the strength is improved at a cost of fracture toughness [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e]. In the case of Mg-ceramic composites (in bulk form), their fracture toughness normally cannot even reach 10 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e when the specific flexural strength exceeds 150 MPa/(g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e); toughnesses above 10 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e have only been attained in lower-strength composites [\u003cspan additionalcitationids=\"CR17\" citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eNatural nacre, which features a layered structure comprising\u0026thinsp;~\u0026thinsp;95 vol.% aragonite platelets with ~\u0026thinsp;5 vol.% organic matter, exhibits much more superior combination of strength and fracture toughness than its constituents [\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e, \u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]. Taking the inspiration from nacre offers an effective approach for developing new high-performance composite materials [\u003cspan additionalcitationids=\"CR22\" citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e]. The so-called lamellar and brick-and-mortar architectures are two of the most common structural arrangements mimicking nacre [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan additionalcitationids=\"CR25 CR26 CR27 CR28\" citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e]. The stiff and soft constituents are alternately arranged in a layered fashion for the former; whereas the stiff platelets (bricks) are more closely stacked and staggered between each other within the soft matrix (mortar) for the latter. These two architectures also show a large difference in their phase constitution, \u003cem\u003ei.e.\u003c/em\u003e, the stiff constituent accounts for a larger volume fraction in the brick-and-mortar architecture than in the lamellar one. Such architectural designs have been implemented in a variety of composite materials, such as those of ceramic-polymer [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e, \u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e], metal-ceramic [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e, \u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e] and metal-graphene systems [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e, \u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e], and been proven to be effective in toughening materials.\u003c/p\u003e \u003cp\u003eHowever, it is not easy to construct nacre-like architectures, especially the brick-and-mortar architectures, in Mg-based composites. This is largely due to the high reactivity of the Mg melt which tends to react at high temperatures with many of the possible reinforcement phases to form brittle intermetallics, or otherwise display a poor wettability with them [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e]. Nevertheless, the mechanical properties of nacre-like composites are closely associated with the type of architectures and the characteristics of matrices [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan additionalcitationids=\"CR25\" citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e], but these factors and their effects remain to be explored for Mg-based composites.\u003c/p\u003e"},{"header":"2. Results","content":"\u003cp\u003eFigure \u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e shows the microstructures and phase constitution of the infiltrated Mg-MAX phase composites. The constituents that were respectively rich in metal and ceramic phases showed an alternate arrangement in a layer fashion in the composites (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea,b), which is qualitatively similar to the structure of natural nacre. The ceramic-rich constituents were interconnected between layers, via bridges for the lamellar architecture and by overlapping for the brick-and-mortar one. As compared to the lamellar architecture, the ceramic-rich constituents in the brick-and-mortar architecture were more closely stacked and accounted for a larger volume fraction in the composite. In both architectures, the ceramic-rich constituents were indeed composed of both metal and ceramic phases instead of monolithic ceramics (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec). The metal phase within these constituents had ultrafine grains typically smaller than 2 \u0026micro;m (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed), due to the constraint of ceramic phase on the growth of metal grains during the solidification process after the melt infiltration procedure.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe ceramic contents in the composites with pure Mg and AZ91D alloy as the metal component were calculated to be ~\u0026thinsp;11.4 vol.% and ~\u0026thinsp;14.0 vol.% for the lamellar architecture, and ~\u0026thinsp;32.6 vol.% and ~\u0026thinsp;32.5 vol.% for the brick-and-mortar architecture, respectively, according to the rule-of-mixtures based on their measured densities (respectively, ~\u0026thinsp;2.02 g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e and ~\u0026thinsp;2.16 g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e for the lamellar architecture, and ~\u0026thinsp;2.54 g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e and ~\u0026thinsp;2.60 g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e for the brick-and-mortar architecture). The volume fractions of the ceramic-rich constituents were determined by image analysis to be ~\u0026thinsp;31.1% and ~\u0026thinsp;51.3% for the pure Mg composites with lamellar and brick-and-mortar architectures, respectively; the corresponding values for the AZ91D alloy composites were ~\u0026thinsp;30.6% and ~\u0026thinsp;52.0%. As such, the volume fractions of the metal phase within the ceramic-rich constituents were derived to be roughly 63.5% and 54.2% for the pure Mg and AZ91D alloy composites with the lamellar architecture, and to be 36.5% and 37.4% for those with the brick-and-mortar architecture, respectively.\u003c/p\u003e \u003cp\u003eEDS analysis revealed that the Mg element and the elemental Ti, Al and C were principally concentrated in the metal and ceramic-rich constituents (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ee and Figure S1 in Supplementary Materials), respectively. This, combined with the XRD results (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ef), clearly indicates that the composites largely retained the pre-designed phase constitution of Mg and Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e phases after melt infiltration. In addition, a small amount of MgO and Mg\u003csub\u003e17\u003c/sub\u003eAl\u003csub\u003e12\u003c/sub\u003e phases were also recognized in the XRD patterns. The former should be the product of the reaction between the Mg melt and the surface oxides of ceramic phase during the melt infiltration process (the oxygen content in the ball-milled Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e platelets was measured to be ~\u0026thinsp;4.2 wt.%). The Mg\u003csub\u003e17\u003c/sub\u003eAl\u003csub\u003e12\u003c/sub\u003e phase may have formed due to the interfacial reaction between the melting Mg and Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e phase [\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e, \u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. Similar reactions have been identified between other metals and MAX phases [\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e, \u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e], \u003cem\u003ee.g.\u003c/em\u003e, in Ag-Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e system [\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e], and been demonstrated to play a role in promoting a strong interfacial bonding between the metal and ceramic phases.\u003c/p\u003e \u003cp\u003eThe nacre-like Mg-MAX phase composites exhibit large differences in their local mechanical properties. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea, for the lamellar architecture subject to an equal load under nanoindentation condition, relatively larger displacements were obtained: (i) in the metal constituent (inter-lamellar region) than in the ceramic-rich constituent (lamella) for the same type of metal component, or (ii) in the composites with pure Mg than AZ91D alloy than in the metal component for the same regions. Specifically, the nanoindentation hardness of the ceramic-rich constituents was markedly higher than that of the inter-lamellar region respectively by ~\u0026thinsp;1.7 and ~\u0026thinsp;2.1 times in the pure Mg and AZ91D alloy composites (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb). Additionally, the lamella and inter-lamellar regions of the AZ91D alloy composites were harder by ~\u0026thinsp;1.4 and ~\u0026thinsp;1.1 times, respectively, than in the pure Mg composites. However, the local elastic moduli determined by nanoindentation testing were similar between the composites containing the two types of metal component, both measured to be ~\u0026thinsp;91.4 GPa and ~\u0026thinsp;61.9 GPa, respectively, for the lamella and inter-lamellar regions.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe global mechanical properties of the composites are also related to the types of nacre-like architecture and metal component. For the same type of metal component, the brick-and-mortar architecture exhibited higher flexural strengths than the lamellar one, respectively up to 491.1\u0026thinsp;\u0026plusmn;\u0026thinsp;27.4 MPa and 652.4\u0026thinsp;\u0026plusmn;\u0026thinsp;45.9 MPa for the pure Mg and AZ91D alloy composites (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec). Nevertheless, this was accompanied by an obvious compromise in the ductility for both metal components. With respect to the effects of metal component, the AZ91D alloy composites invariably displayed relatively higher flexural strengths when considering the same type of architecture, whereas the pure Mg composites displayed better ductility. All these composites showed microscopically staircase-like morphologies on their fracture surfaces, which are consistent with their layered structure (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed-f). Apparent dimples that are indicative of plastic deformation were formed in the metal constituent for the pure Mg composites with lamellar architecture (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed,e). Such feature became much less evident in the AZ91D alloy composites (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ef), consistent with their reduced ductility.\u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea shows the representative single-edge notched bending load-displacement curves of the nacre-like Mg-MAX phase composites. Instantaneous fracture occurred at the maximum bending load without stable crack propagation in the AZ91D alloy composites with the brick-and-mortar architecture. In comparison, the bending load decreased with increasing displacement after reaching its maximum in the other composites. In particular, the pure Mg composites with the lamellar architecture showed gradual load drop along with a large displacement, indicating a stable cracking process. The fracture resistance of the composites in terms of \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J\\)\u003c/span\u003e\u003c/span\u003e-integral and equivalent stress intensity \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({K}_{J}\\)\u003c/span\u003e\u003c/span\u003e as a function of crack extension (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\varDelta a\\)\u003c/span\u003e\u003c/span\u003e) are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb,c. The AZ91D alloy composites with brick-and-mortar architecture fractured abruptly at the onset of crack propagation, giving \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(K\\)\u003c/span\u003e\u003c/span\u003e of 2.1\u0026thinsp;\u0026plusmn;\u0026thinsp;0.7 kJ\u0026middot;m\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e and 14.4\u0026thinsp;\u0026plusmn;\u0026thinsp;2.7 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e, respectively. However, all the other composites, \u003cem\u003ei.e.\u003c/em\u003e, the AZ91D alloy composites with lamellar architecture and the pure Mg composites, displayed a stable increase in \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({K}_{J}\\)\u003c/span\u003e\u003c/span\u003e with crack extension, showing rising crack resistance-curve (R-curve) behavior. Higher fracture toughness was generated in the lamellar architecture for both types of metal component, or in the pure Mg composites for a given type of architecture. In particular, the pure Mg composites with lamellar architecture exhibited the highest fracture toughness, with the critical \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J\\)\u003c/span\u003e\u003c/span\u003e and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({K}_{J}\\)\u003c/span\u003e\u003c/span\u003e up to respectively 15.9\u0026thinsp;\u0026plusmn;\u0026thinsp;2.5 kJ\u0026middot;m\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e and 32.5\u0026thinsp;\u0026plusmn;\u0026thinsp;2.6 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e, which were determined by the maximum crack extension of ~\u0026thinsp;0.5 mm according to the ASTM Standard E1820 (as indicated by the dashed lines) [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe nacre-like architectures have been proven to be effective in activating a series of extrinsic toughening mechanisms for toughening materials, such as crack deflection, uncracked-ligament bridging and frictional sliding between crack faces [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e]. These mechanisms were clearly introduced into the Mg-MAX phase composites by the lamellar architecture, as manifested by the obviously tortuous cracking paths (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed); nevertheless, such behavior became much less evident for the brick-and-mortar architecture. With regard to the effects of the type of metal component, obvious plastic deformation occurred in the inter-lamellar region near the primary crack in the pure Mg composites with lamellar architecture; whereas microcracks were formed in the ceramic-rich constituents (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee). In comparison, for the AZ91D alloy composites, both lamella and inter-lamellar regions exhibited apparent microcracking without visible plastic deformation (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ef).\u003c/p\u003e"},{"header":"3. Discussion","content":"\u003cp\u003eBoth plastic deformation and microcracking in the vicinity of the crack tip, which were dominant respectively in pure Mg and AZ91D alloy composites, offer a means for dissipating mechanical energy and can promote crack bifurcation (Figure S2). These mechanisms, which can be termed intrinsic toughening mechanisms [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e], are principally active ahead of the crack tip to blunt the crack tip. In addition, they play a role in promoting the crack deflection following the slip bands in inter-lamellar regions or along the microcracks (Figure S2) and inducing the formation of uncracked-ligament bridges, \u003cem\u003ei.e.\u003c/em\u003e, between the main cracks and voids or microcracks. These mechanisms, which can be termed extrinsic toughening mechanisms [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e], are principally active at or behind the crack tip to reduce the crack-driving force actually experienced at the crack tip, thereby shielding the crack tip from applied stress.\u003c/p\u003e \u003cp\u003eThe effectiveness of the above toughening mechanisms in the nacre-like Mg-MAX phase composites is associated with the types of architectures (lamellar or brick-and-mortar) and metal components (pure Mg or AZ91D alloy). The lamellar architecture, which contains a larger volume fraction of metal constituent than the brick-and-mortar one, enables a more stable crack extension via crack defection, bifurcation and uncracked-ligament bridging mechanisms, leading to higher crack-growth toughness after the onset of cracking. Such trend is opposite to that found in nacre-like composites of ceramic-polymer (polymethylmethacrylate or epoxy resin) [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e] and alumina-metallic glass [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e] systems where the brick-and-mortar architecture is tougher than the lamellar one. Indeed, the relatively soft constituents (of polymers or metallic glass) that comprise the \u0026ldquo;mortar\u0026rdquo; appear to be brittle and can hardly undergo large plastic deformation, which is different from the case for pure Mg and AZ91D alloy. As such, the fracture resistance in these composites is mainly developed from the extrinsic toughening mechanisms which are more effective in the brick-and-mortar architecture owing to its higher density of interfaces and bridges. With respect to the effects of the metal component, pure Mg exhibits a lower strength than the AZ91D alloy in the as-cast state (Figure S3), and thereby can be more easily deformed in the Mg-MAX phase composites. Such easy inelastic deformation of the soft phase, as compared to the fracture of stiff constituents, is essential for ensuring the toughening effects of the nacre-like architectures in promoting both the intrinsic and extrinsic mechanisms [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea presents a comparison of the fracture toughness and flexural strength for the Mg-MAX phase composites with their component materials of pure Mg, AZ91 (Mg-Al-Zn) series alloys and bulk Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e ceramics [\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e, \u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e, \u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e, \u003cspan additionalcitationids=\"CR38 CR39\" citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e]. Most MAX phase ceramics were reported to exhibit a flexural strength of less than 560 MPa (although they can be strengthened by fine modulation of their microstructures [\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e]); and their fracture toughness rarely exceeds 10 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e [\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e]. Pure Mg and AZ91 series alloys generally exhibit a fracture toughness of lower than 24 MPa\u0026middot;m\u003csup\u003e1/2\u003c/sup\u003e and a flexural strength below 400 MPa [\u003cspan additionalcitationids=\"CR38 CR39\" citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e]. In comparison, the flexural strength of the nacre-like Mg-MAX phase composites is comparable to, or even higher than, those of monolithic Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e ceramics, while their fracture toughness is comparable to that of the metal phase. This indicates the achievement of a good combination of property advantages in the composites inherited from their constituent materials. Figure\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb shows the variation in the flexural strength as a function of the ceramic content in Mg-based composites reinforced by different types of ceramics [\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e, \u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e, \u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e, \u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e, \u003cspan additionalcitationids=\"CR43\" citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e]. The current Mg-MAX phase composites exhibit relatively higher flexural strength than most other materials at equal ceramic contents. Figure\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec shows a comparison of the specific fracture toughness and specific flexural strength, \u003cem\u003ei.e.\u003c/em\u003e, fracture toughness and flexural strength normalized by density, for the nacre-like composites of different metal-ceramic systems fabricated by ice-templating and subsequent melt infiltration procedures [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e, \u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e, \u003cspan additionalcitationids=\"CR46 CR47 CR48\" citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e]. The Mg-MAX phase composites show a high specific flexural strength ranging from ~\u0026thinsp;200 MPa/(g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e) to over 250 MPa/(g\u0026middot;cm\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e) combined with a high specific fracture toughness. In particular, their specific fracture toughness exceeds those of most other nacre-like metal-ceramic composites with the same level of specific flexural strength. This good combination of properties makes the nacre-like Mg-MAX phase composites appealing for many lightweight structural applications.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eIn summary, a set of Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures were fabricated by pressureless infiltration of pure Mg and AZ91D alloy melt into ice-templated porous Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e scaffolds. The ceramic-rich constituents were alternately arranged with the metals in a layered fashion in the composites, and were more densely stacked in the brick-and-mortar architecture than in the lamellar one. They had an ultrafine structure comprising both ceramic and metal phases that were interconnected between adjacent layers. The mechanical properties of the Mg-MAX phase composites were associated with the types of architectures and metal components. Specifically, higher flexural strength, yet accompanied with inferior ductility, was generated in the AZ91D (Mg-Al-Zn-Mn) magnesium alloy composites for a given architecture or in the brick-and-mortar architecture for a given metal component. Apart from the AZ91D alloy composites with brick-and-mortar architecture, the Mg-MAX phase composites displayed rising R-curve behavior with stable crack propagation process. Crack advance was seen to involve bifurcation with crack blunting via plastic deformation and microcracking, which were respectively dominant in pure Mg and AZ91D alloy composites; additionally, these materials were extrinsically toughened by crack deflection and uncracked-ligament bridging resulting in significant crack-tip shielding, although these latter mechanisms were more evident in the lamellar than in the brick-and-mortar architectures. The composites achieved an effective combination of the property advantages from their component materials. They were stronger than most other Mg-ceramic composites at equal ceramic contents and exhibited a good combination of specific fracture toughness and specific flexural strength among nacre-like metal-ceramic composites, thereby demonstrating a potential for lightweight structural applications. We believe that his study may further give insights for the design of new high-performance Mg-based composites.\u003c/p\u003e"},{"header":"4. Methods","content":"\u003cp\u003eCommercial Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e powders (400 mesh, Haixin metal, China) were ball milled in an ethanol dispersant in a mixed gas of 20 vol.% oxygen and 80 vol.% argon at a speed of 500 rpm for 48 h. In this process, the MAX phase powders were delaminated into ultrafine platelets along their relatively weak MX/A interlayer bonding of layered structure [\u003cspan citationid=\"CR50\" class=\"CitationRef\"\u003e50\u003c/span\u003e]. The equivalent particle size of the platelets was measured to be 100\u0026ndash;300 nm using laser diffraction granulometry. The platelets were dried in air for 2 days and then were dispersed in deionized water at a mass ratio of 1:2 under ultrasonic treatment at 50 W for 0.5 h. The slurry was mixed with an anionic dispersant (Darvan CN, R.T. Vanderbilt, USA), polyvinyl alcohol (Meryer, China) and hydroxypropyl methylcellulose (Meryer, China), which accounted, respectively, for 0.2 wt.%, 1.5 wt.% and 0.5 wt.% of the platelets. These additives were used for promoting dispersion, increasing the viscosity of slurry to mitigate gravitational sedimentation, and binding the platelets after the removal of the water.\u003c/p\u003e \u003cp\u003eThe slurry was ball milled at a speed of 25 rpm for 24 h, and then was poured into a cuboid Teflon mold with an inner square cross-section of 30 mm \u0026times; 30 mm which was fixed onto a customized freezing apparatus. The mold was designed to have non-uniform lateral walls which was thinner at one side (3 mm) than the other three sides (20 mm), and was sealed using a wedge with a slope angle of 25\u0026deg; at its bottom with the thinner end of the wedge conforming to the thinner wall of the mold. Such designs were implemented to create a horizontal temperature gradient in addition to the vertical one in the slurry during the freezing process [\u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e, \u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e]. The mold was descended at a constant rate of 0.6 mm\u0026middot;min\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e to immerse into an insulated box containing liquid nitrogen until the slurry was fully frozen. The frozen body was demolded and freeze-dried in a vacuum below 5 Pa for at least 72 h using a Scientz-10ND freeze drier (Scientz Biotechnology, China). Porous Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e scaffolds with lamellar architecture were obtained by sintering the freeze-dried samples in flowing argon gas at 900\u0026deg;C for 1 h. Some of the scaffolds were infiltrated with paraffin wax and then uniaxially pressed along the normal direction of lamellae at 75\u0026deg;C (~\u0026thinsp;15\u0026deg;C higher than the melting point of paraffin wax) under a pressure of 20\u0026thinsp;~\u0026thinsp;35 MPa for 2 h for densification. These scaffolds were re-sintered in flowing argon gas at 950\u0026deg;C for 1 h to create interconnections between adjacent lamellae or bricks.\u003c/p\u003e \u003cp\u003ePure Mg with a purity of 99.99 wt.% and a widely used AZ91D cast Mg alloy (with ~\u0026thinsp;8.5\u0026ndash;9.5 wt.% Al, ~\u0026thinsp;0.45\u0026ndash;0.9 wt.% Zn and ~\u0026thinsp;0.17\u0026ndash;0.5 wt.% Mn) were used for fabricating the Mg-MAX phase composites. The sintered Ti\u003csub\u003e3\u003c/sub\u003eAlC\u003csub\u003e2\u003c/sub\u003e scaffolds were infiltrated with the melt of pure Mg or AZ91D alloy in flowing argon gas at 850\u0026deg;C, some 250\u0026deg;C higher than their melting points, for 1.5 h. The densities of the infiltrated composites were measured using the Archimedes\u0026rsquo; method [\u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e]. The phase constitution was determined by X-ray diffraction (XRD) using a Bruker D8 Advance X-ray diffractometer (Bruker AXS, Germany) with Cu-Kα radiation. Scanning electron microscopy (SEM) imaging was conducted using an Inspect F50 field-emission scanning electron microscope (FEI, USA) operating at an accelerating voltage of 20 kV. Energy dispersive X-ray spectroscopy (EDS) measurements were carried out using a Bruker EDS system attached to the microscope. Scanning transmission electron microscopy (STEM) imaging was performed using a Tacnai-G2 F30 transmission electron microscope (FEI, USA) operating at an accelerating voltage of 300 kV.\u003c/p\u003e \u003cp\u003eNanoindentation tests were performed using a G200 nano indenter (KLA, USA) with a Berkovich diamond tip at a constant loading rate of 0.67 mN\u0026middot;s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e to a peak load of 10 mN with a holding time of 10 s followed by unloading. Three-point bending tests were conducted using an Instron E1000 testing system (Instron, USA) at room temperature with a constant displacement rate of 0.1 mm\u0026middot;min\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e on beam samples with dimensions of 2 mm \u0026times; 1.5 mm \u0026times; 25 mm and a loading span of 20 mm, in line with the ASTM Standard C1161 [\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e]. Single-edge notched bending tests were performed on a JEOL MicroTest stage inside the chamber of a JEOL JSM-6510 scanning electron microscope (JEOL, Japan) at a displacement rate of 5 \u0026micro;m\u0026middot;min\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. Samples with a width of 4 mm, thickness of 2 mm and loading span of 16 mm were side notched to ~\u0026thinsp;2 mm in depth using a low-speed diamond wire saw with the notch sharpened to a tip radius of 5\u0026thinsp;~\u0026thinsp;10 \u0026micro;m using a razor blade with 1 \u0026micro;m diamond paste, in general accordance with the ASTM Standard E1820 [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e]. The loads in both three-point and singe-edge notched bending tests were perpendicular to the layered structure of the composites, which is consistent with the general loading configuration for natural nacre. The fracture morphologies of samples were characterized by SEM imaging.\u003c/p\u003e \u003cp\u003eNonlinear-elastic fracture mechanics methods were used to evaluate the fracture toughness of the composites in terms of their \u003cem\u003eJ\u003c/em\u003e-integral. The \u003cem\u003eJ\u003c/em\u003e values were calculated from the applied load and instantaneous crack length by considering the elastic (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{el}\\)\u003c/span\u003e\u003c/span\u003e) and plastic (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{pl}\\)\u003c/span\u003e\u003c/span\u003e) components, \u003cem\u003ei.e.\u003c/em\u003e, \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J={J}_{el}+{J}_{pl}\\)\u003c/span\u003e\u003c/span\u003e, according to the ASTM Standard E1820 [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e]. \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{el}\\)\u003c/span\u003e\u003c/span\u003e was formulated using the linear-elastic stress intensity factor \u003cem\u003eK\u003c/em\u003e according to \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{el}=\\frac{{K}^{2}\\left(1-{v}^{2}\\right)}{E}\\)\u003c/span\u003e\u003c/span\u003e, where \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(E\\)\u003c/span\u003e\u003c/span\u003e is the Young\u0026rsquo;s modulus and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(v\\)\u003c/span\u003e\u003c/span\u003e is the Poisson\u0026rsquo;s ratio. The \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(E\\)\u003c/span\u003e\u003c/span\u003e values were approximated using the averages of those calculated following the rule-of-mixtures under the iso-strain (Voigt model) and iso-stress (Reuss model) conditions [\u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e]. In view that the mode I stresses were parallel to the layered structure of the nacre-like composites, such approach was expected to give a conservative approximation for the Young\u0026rsquo;s modulus and therefore the fracture toughness. The Poisson\u0026rsquo;s ratio \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(v\\)\u003c/span\u003e\u003c/span\u003e of the composites was approximated in the same fashion to be 0.28 and 0.26, respectively, for the lamellar and brick-and-mortar architectures. The plastic component \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{pl}\\)\u003c/span\u003e\u003c/span\u003e was calculated from the plastic area (\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({A}_{pl}\\)\u003c/span\u003e\u003c/span\u003e) under the load-displacement curve according to \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({J}_{pl}=\\frac{1.9{A}_{pl}}{Bb}\\)\u003c/span\u003e\u003c/span\u003e [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e], where \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(B\\)\u003c/span\u003e\u003c/span\u003e is the thickness of sample and \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(b\\)\u003c/span\u003e\u003c/span\u003e is the length of the uncracked ligament. The equivalent \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(K\\)\u003c/span\u003e\u003c/span\u003e-based stress intensity was obtained from the \u003cem\u003eJ\u003c/em\u003e-integral based on the standard \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(J\\)\u003c/span\u003e\u003c/span\u003e-\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(K\\)\u003c/span\u003e\u003c/span\u003e equivalence as \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\({K}_{J}={\\left(\\frac{JE}{\\left(1-{v}^{2}\\right)}\\right)}^{1/2}\\)\u003c/span\u003e\u003c/span\u003e in the case of mode I fracture [\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e].\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe data that support the findings of this study are available from the corresponding author, Prof. Zengqian Liu, at [email protected], upon reasonable request.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAcknowledgements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors are grateful for the financial support by the\u0026nbsp;National Key R\u0026amp;D Program of China (2020YFA0710404),\u0026nbsp;the National Natural Science Foundation of China (52173269\u0026nbsp;and\u0026nbsp;51871216), the KC Wong Education Foundation (GJTD-2020-09),\u0026nbsp;the Youth Innovation Promotion Association CAS, and the CAS Project for Young Scientists in Basic Research (YSBR-025). ROR was supported by the Multi-University Research Initiative (AFOSR-FA9550-15-1-0009) from the Air Force Office of Scientific Research.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eDeclaration of competing interest\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare no conflict of interest.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003ePollock TM. Weight loss with magnesium alloys. Science 2010;328:986-987.\u003c/li\u003e\n\u003cli\u003eChen LY, Xu JQ, Choi HS, Pozuelo M, Ma XL, Bhowmick S, Yang JM, Mathaudhu S, Li XC. Processing and properties of magnesium containing a dense uniform dispersion of nanoparticles. Nature 2015;528:539-543.\u003c/li\u003e\n\u003cli\u003eNguyen QB, Tun KS, Lim CYH, Wong WLE, Gupta M. Influence of nano-alumina and submicron copper on mechanical properties of magnesium alloy AZ31. Compos Part B 2013;55:486-491.\u003c/li\u003e\n\u003cli\u003eWang Y, Wang HY, Xiu K, Wang HY, Jiang QC. 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Continuous ice-templating of macro-porous materials with uniformly ordered architecture. Sci China Mater. doi:10.1007/s40843-022-2078-8.\u003c/li\u003e\n\u003cli\u003eHughen SW. Archimedes revisited: a faster, better, cheaper method of accurately measuring the volume of small objects. Phys Educ 2005;40:468-474.\u003c/li\u003e\n\u003cli\u003eASTM Standard C1161-18, \u003cem\u003eStandard Test Method for Flexural Strength of Advanced Ceramics at Ambient Temperature\u003c/em\u003e, American Society for Testing and Materials International, West Conshohocken, PA, USA 2018.\u003c/li\u003e\n\u003cli\u003eWatt JP, Davies GF, O\u0026rsquo;Connell RJ. The elastic properties of composite materials. Rev Geophys Space Phys 1976;14:541-563.\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"Magnesium composites, Bioinspired designs, Nacre-like architectures, Fracture toughness, MAX phase","lastPublishedDoi":"10.21203/rs.3.rs-1896279/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-1896279/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eBioinspired nacre-like structures are effective in toughening materials, yet are difficult to construct in Mg-ceramic systems. Here, a set of Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures were fabricated by pressureless infiltration of the Mg melt into ice-templated ceramic scaffolds. The structure and mechanical properties of the composites were elucidated with a special focus on the effects of the types of architectures (lamellar or brick-and-mortar) and matrices (pure Mg or Mg alloy) on the toughening mechanisms. The nacre-like architectures were found to play a role in blunting the cracks via plastic deformation and microcracking, and shielding the cracks from applied stress by promoting crack deflection and uncracked-ligament bridging mechanisms. These composites achieved a good combination of (specific) strength and fracture toughness which are superior to other reported Mg-ceramic and nacre-like metal-ceramic composite materials.\u003c/p\u003e","manuscriptTitle":"Strong and tough Mg-MAX phase composites with nacre-like lamellar and brick-and-mortar architectures","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2022-07-29 18:55:42","doi":"10.21203/rs.3.rs-1896279/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"communications-materials","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"commsmat","sideBox":"Learn more about [Communications Materials](https://www.nature.com/commsmat/)","snPcode":"","submissionUrl":"","title":"Communications Materials","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Communications Series","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"2f9701e2-2954-4ac6-8a24-fd3d34191255","owner":[],"postedDate":"July 29th, 2022","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[],"tags":[],"updatedAt":"2023-05-11T21:00:59+00:00","versionOfRecord":{"articleIdentity":"rs-1896279","link":"https://doi.org/10.1038/s43246-023-00358-3","journal":{"identity":"communications-materials","isVorOnly":false,"title":"Communications Materials"},"publishedOn":"2023-05-09 04:00:00","publishedOnDateReadable":"May 9th, 2023"},"versionCreatedAt":"2022-07-29 18:55:42","video":"","vorDoi":"10.1038/s43246-023-00358-3","vorDoiUrl":"https://doi.org/10.1038/s43246-023-00358-3","workflowStages":[]},"version":"v1","identity":"rs-1896279","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-1896279","identity":"rs-1896279","version":["v1"]},"buildId":"7rjqhiLT3MXkJMwkYKINL","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}