aluminum matrix composites

Journal Pre-proof Enhancement of strength and ductility by interfacial nano-decoration in carbon nanotube/aluminum matrix composites Baisong Guo, Yiqiang Chen, Zhangwei Wang, Jianhong Yi, Song Ni, Yong Du, Wei Li, Min Song PII:

S0008-6223(19)31273-4

DOI:

https://doi.org/10.1016/j.carbon.2019.12.038

Reference:

CARBON 14888

To appear in:

Carbon

Received Date: 18 September 2019 Revised Date:

16 December 2019

Accepted Date: 16 December 2019

Please cite this article as: B. Guo, Y. Chen, Z. Wang, J. Yi, S. Ni, Y. Du, W. Li, M. Song, Enhancement of strength and ductility by interfacial nano-decoration in carbon nanotube/aluminum matrix composites, Carbon (2020), doi: https://doi.org/10.1016/j.carbon.2019.12.038. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Ltd.

Author contributions Section M.S., Z.W. and J.Y. conceived the study. B.G. and Y.C. performed the experiments (under supervision by M.S., S.N. and W.L.). B.G., M.S., Z.W. and Y.D. analyzed the results and wrote the manuscript. All authors participated in the discussion and interpretation of the results.

Enhancement of strength and ductility by interfacial nano-decoration in carbon nanotube/aluminum matrix composites Baisong Guo1,2‡, Yiqiang Chen3‡, Zhangwei Wang3**, Jianhong Yi4, Song Ni1, Yong Du1, Wei Li2, Min Song1* 1. State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, China 2. Institute of Advanced Wear & Corrosion Resistant and Functional Materials, Jinan University, Guangzhou, 510632, China 3. Max-Planck-Institut für Eisenforschung, Max-Planck-Str. 1, 40237 Düsseldorf, Germany. 4. School of Materials Science and Engineering, Kunming University of Science and Technology, Kunming, 650093, China

Abstract Aluminum (Al) matrix composites gain tremendous attention as candidates for lightweight structural materials. Interfaces between the matrix and reinforcements, long-standing concerns, are critical in determining the mechanical properties of Al matrix composites. Unlike the conventional thoughts that focus on raising the interfacial wettability, a novel interfacial nanodecoration strategy is reported to enhance the interfacial adhesion by forming a diffusion interface between Al and carbon nanotubes (CNTs) through copper (Cu) coating on the surface of CNTs. The resulted Cu-rich nanolayers through this strategy compromise the large interfacial misfit strain between Al and CNTs. Such unique interfacial structure improves the strengthening efficiency of CNTs and benefits the plastic deformation of the Al matrix, and thus contributes to a simultaneous increase in strength and ductility and breaks the ubiquitous strength-ductility trade-off dilemma in the structural material design. Consequently, we achieve an exceptional combination of tensile strength (391 MPa) and tensile elongation (15.7 %) for our composite that surpasses its counterparts. The present tactic thus paves a new way to process high-performance Al matrix composites.

*Corresponding author. E-mail: [email protected] (M. S.), Tel.: +86(0)73188877677 **Corresponding author. E-mail: [email protected] (Z. W.), Tel.: +49(0)2116792275 ‡ These authors contributed equally to this work.

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1. Introduction Aluminum (Al) matrix composites are important lightweight structural materials that have been widely applied in industries of aircraft, aerospace and automobile [1-3]. Over the past two decades, carbon nanotubes (CNTs) have been regarded as one of the most intriguing reinforcements for Al matrix composites, since the high Young's modulus (~270-950 GPa) [4], high tensile strength (>110 GPa) [5], and low density (~0.5-2.8 g/cm3) [6] of CNTs outcompete those of the traditional ceramic particles. However, the high potential applicability of CNTs to reinforcing Al matrix was hindered by the difficulty in uniformly dispersing CNTs into the Al matrix and also the weak interface between Al and CNTs. To overcome the dispersion problem, several methods such as in-situ growth [7], flake powder metallurgy [8-10], solution ball milling [11], high energy ball milling [12-15] and molecular level mixing [16, 17] have been developed. For example, Esawi et al. [12] employed the high energy ball milling method to achieve the uniform dispersion of 2 wt. % CNTs in the Al matrix, resulting in enhancement of around 21 % in tensile strength compared to the equivalent of pure aluminum. Jiang et al. [18] produced uniformly dispersed 1 vol. % CNTs reinforced Al matrix composites, which increases the tensile strength from 330 MPa to 375 MPa. Nevertheless, the strengthening efficiency of CNTs is still far less than our expectation. Moreover, the obtained increases of strength in most of studies are accompanied by the substantially deteriorated ductility. In addition to the dispersion condition of CNTs in the Al matrix, it is well established that the interfacial adhesion between the matrix and reinforcements plays a vital role in controlling the mechanical properties of Al matrix composites [19, 20]. The weak interfaces lead to low strengthening effects of reinforcements [20]. Upon loading, the stress transfer from the soft matrix to the hard reinforcments is one of the main strengheing mechanims in composites. The weakly bonded interfaces limit the stress transfer and further hinder the pursuit of high mechanical performance of the composites. In addition, these weakly bonded interfaces are prone to nucleate cracks upon loading, which serves as detrimental sources at operations in harsh environments [21]. Therefore, strategies to improve interfaces have been rising to prominence, given the aim of developing Al composites with high-strength and large ductility. By far, most of solutions for strong interfaces between the Al matrix and CNTs concentrate on interfacial reaction [22, 23] and ceramic coating [24-26]. For example, Zhou and Chen et al.

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[22, 23] pointed out that the chemical reaction between Al and CNTs can be beneficial for load transfer from the Al matrix to CNTs and thus hinders the interface slippage, leading to improved strength and good ductility. However, the critical drawback of this tactic is that the limited interfacial reaction is difficult to realize under high temperature condition, because the nanosized CNTs are easy to be completely consumed, leading to the formation of Al4C3 rods [27]. The excess formation of hydrolysable Al4C3 phase would deteriorate the mechanical property and physical/chemical stability of Al/CNTs composites. Saba and Zhang et al. [24, 25] produced TiC or SiC coating on the surface of CNTs to improve the interfacial wetting, and thus to build more strongly bonded interface between Al and CNTs. However, the role of wettability in strengthening the interfacial bonding would be weakened when the Al/CNTs composites were fabricated using the most popular solid state processing, mainly powder metallurgy technique. So, the ideal interfacial bonding should be limited diffusion bonding or chemical bonding, which can largely improve the load transfer efficiency but not destroy the CNTs structure. In this work, an innovative interfacial nano-decoration approach is proposed to enhance the interfacial bonding by forming Cu rich nanolayers between CNTs and the Al matrix, which raises both strength and ductility in CNTs reinforced pure Al composites. The reported composite exhibits a superior combination of strength (391 MPa) and tensile elongation (15.7 %) that exceeds currently existing ones. Our strategy offers valuable insights into producing metal matrix composites with outstanding strength and ductility.

2. Experimental 2.1 Nano-decoration on CNTs The pristine CNTs with an average diameter of 10 nm and length of 3 µm were purchased from Sigma-Aldrich Co. LLC, which were manufactured by the catalytically chemical vapor deposition (CVD) method. Prior to electroless plating, CNTs were pre-treated based on following steps: (i) 1 g raw CNTs (99.0 % in purity) were purified and oxidized in the chemical solution (a mixture of 68 wt. % HNO3 and 95 wt. % H2SO4 with a volume ratio of 3:1) at 90°C for 2 h, with the aid of magnetic stirring; (ii) the acid treated CNTs were further sensitized in the 800 ml aqueous solution (20 ml 37 wt. % HCl with the addition of 22.6 g SnCl2) for 40 minutes at room temperature; (iii) the sensitized CNTs were immersed in 1000 ml activating aqueous solution with 4.59 g AgNO3 for 20 min at room temperature, and then 25 wt. % ammonium

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hydroxide was added into the activating solution until the existing sediments were completely dissolved. After pre-treatments to generate the nucleation sites on the surface of CNTs, the metallization assisted with mechanical stirring was used to coat Cu nanoparticles on CNTs in a plating bath [28-30], followed by dilution via deionized water until reaching neutral pH. The details in the chemical composition and operation condition for the electroless platting solution are listed in Table 1.

Table 1 Compositions of electroless plating solutions. Component

Content

Function

CuSO4·5H2O

12.1 g

Cu ion source

C6H5Na3O7·2H2O

15.5 g

Complexing agent

HCHO (37%)

100 ml

Reducing agent

H2NCSNH2 (＞99%)

100 ml

Stabilizing agent

Deionized water

1000 ml

Solvent

Temperature

80 °C

-

PH (adjust with NaOH)

~9.3

~

2.2 Fabrication of composites Spherical pure Al powders (99.9 % in purity) with average diameter of 2 µm were firstly pre-milled to flakes with a size of ~30 µm. After that, the Cu coated CNTs (2.12 g) and uncoated CNTs (1 g) were mixed with 127 g pre-milled Al powders, respectively, using a planetary ball miller (a ball to powder weight ratio of 3:1) with a rotation speed of 300 rpm for 5 h and 150 ml absolute ethyl alcohol as the medium. In the two types of powder mixtures, the CNTs contents regardless of the existence of Cu are both 1 vol.% by taking the theoretical density of Cu, Al and CNTs as 8.9 g/cm3, 2.7 g/cm3 and 2.1g/cm3, respectively [31]. The chemical compositions of two types of powder mixtures and bulk composites can be denoted as Al-1 vol.%-0.26 vol.% Cu and Al-1 vol.% CNTs. The mixtures were dried thoroughly at 75°C in a vacuum oven, followed by sintering in an FCT HPD 25/3 spark plasma sintering furnace at 600°C for 20 min under argon atmosphere, with an applied pressure of 30 MPa. After sintering, cylinder-shaped specimens with a diameter of 40 mm and height of 15 mm were sealed into stainless steel cans and

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subsequently hot rolled to a height reduction of 50 % through 10 passes at 450°C under the protection of stainless steel sheath. The Cu coated and uncoated CNTs reinforced Al matrix composites were denoted as Al-(nano-Cu)-CNTs and Al-CNTs, respectively. For comparison, pure Al specimens were also prepared using the same sintering and hot-rolling processes.

2.3 Mechanical tests Tensile tests were performed at room temperature using an Instron 3369 testing machine with a strain rate of 2.1×10-3 s-1. The tensile specimens with a diameter of 7 mm and a gauge length of 20 mm were machined by wire-electrode cutting from the rolled sheets along the rolling direction. The average values were acquired from three independent measurements.

2.4 Microstructure characterization Scanning electron microscope (SEM) and transmission electron microscope (TEM observations were performed by an FEI Nova Nano 230 instrument, and a Titan G2 60-300 image-corrected electron microscope, respectively. Z-contrast annular dark field (ADF) imaging was conducted by a probe-corrected FEI Titan Themis 80-300 operating at 300 kV and fitted with an X-FEG high brightness electron source with an elevation angle of ∼18°. A 23.8 mrad probe-forming aperture was used, corresponding to ∼0.8 Å resolution, and an inner collection semi-angle of 73 mrad and an outer collection semi-angle of 200 mrad. The electron backscatter diffraction (EBSD) analysis was performed by using a Philips XL 30S SEM equipped with the TSL-OIM system. An X-ray photoelectron spectroscopy (XPS, Thermo Fisher K-alpha 1083), employing micro-aggregation monochromator as an X-ray source, was applied to detect the structure information of the coated CNTs. For the TEM observation and EBSD analysis, 3 mm thin disks were firstly mechanical grinded to ~100 µm, and further thinned by the Gatan Precision Ion Polishing System operated at 5 kV.

3. Results and discussion 3.1 Microstructures of the composites Fig. 1 illustrates the processing procedure for Al-(nano-Cu)-CNTs composite based on this strategy, mainly including two stages, i.e. nano-decoration on CNTs and composite fabrication. During the nano-decoration process, the oxidization by acid mixture was firstly employed to

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create oxygen containing functional groups on the surface of CNTs, which can enhance the adsorption of Sn2+ [28]. After that, as shown in Fig.1, the Sn2+ is uniformly adsorbed on the surface of CNTs. Subsequently, Sn2+ can react with the [Ag(NH3)2]+, resulting in the formation of Ag(0) particles, which can act as the catalyst to promote the reductive reaction of Cu2+ [32]. When the CNTs after sensitization and activation were immersed into the plating solution, the Cu2+ would be reduced to Cu nanoparticles on the CNTs surface by the reducing agent of HCHO [33]. Furthermore, the Cu2+ in the solution would be adsorbed on the Cu nanoparticles to form bigger Cu particles.

Figure 1 Schematic illustration for the processing route of the Al-(nano-Cu)-CNTs composite.

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Fig. 2 shows the microstructures of the acid treated CNTs before and after coating by Cu nanoparticles. The TEM image in Fig. 2a shows that CNTs retain their typical tubular shapes after a HNO3-H2SO4 acid treatment. The acid treatment removed the impurities in raw CNTs— smooth and clean surfaces are exhibited, as revealed by the high resolution TEM (HRTEM) image in Fig. 2c. After the coating process, numerous nanosized particles can be observed on surfaces of CNTs (see the inset in Fig. 2b), yet the overall length and shape of CNTs change very little by comparing Fig. 2a with Fig. 2b. Fig. 2d shows that the Cu-coated CNTs possess preferable rough surfaces in contrast to the smooth surfaces of uncoated CNTs (see Fig. 2c).

Figure 2 (a) TEM micrograph of the HNO3-H2SO4 acid treated CNTs before Cu coating. (b) TEM micrograph of the acid treated CNTs after Cu coating and the inset shows Cu nanoparticles attached on surfaces of CNTs. (c) and (d) are HRTEM micrographs of CNTs before and after Cu coating, respectively.

The X-ray photoelectron spectroscopy (XPS) analysis was performed for the CNTs after coating process. The XPS full spectrum (Fig.3a) shows four primary XPS peaks of C1s，O1s,

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Cu LM2 and Cu 2p3 detected in the Cu coated CNTs specimen, which reveals that the atomic ratio of Cu and C is 17.41 at.% and 82.59 at.%, respectively. The deconvolution of the XPS C1s, O1s, Cu2p3 peaks are shown in Fig. 3b, and c, d, respectively. A strong peak at 284.0 eV and a weak peak at 285.9 eV corresponding to the sp2 and sp3 of C-C can be clearly observed in Fig.3b [34], which indicate that the original structure of CNTs was maintained after the electroless platting procedure. In addition, the peak at 287.7 eV corresponding to C=O proves that acid treatments created functional groups on the surface of CNTs [35], which enable the subsequent attachment of Cu nanoparticles. Fig.3c shows three O1s peaks, with the fitted peak values at 530.3 eV, 532.1 eV and 535.5 eV, corresponding to the physically absorbed O, C-O or O-C=O, π bonded O, respectively [36]. Fig.3d displays the peaks at 932.5 eV and 934.5 eV corresponding to Cu and CuO [37], which further confirms that most of these attached nanoparticles on the surfaces of CNTs (see Fig. 2b) are Cu. Although the occurrence of CuO peak, it can be believed that the oxidation extent of coated Cu is very slight, because most of detected oxygen comes from the physical adsorption and residual oxygen containing functional groups (see Fig. 3c). Moreover, the O in Cu oxide can be effectively removed during the SPS process [38], and the slight oxidization of Cu coating will not exert any adverse effect on the mechanical properties of bulk Al/CNTs composites.

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Figure 3 Deconvolution of the XPS: (a) Cu2p, (b) C1s, and (b) O1s peaks for the Cu coated CNTs.

Fig. 4 a-d show microstructures of mixtures of Al powders and CNTs. Since the as-received spherical Al powders were pre-milled to flakes before mixing, a shape that favors absorption of CNTs [8], homogeneously dispersed CNTs on the surfaces of Al flakes (see Fig. 4c and d) have been achieved. The recorded X-ray spectra (insets in Fig. 4c and d) based on the SEM-EDS analysis further confirms the existence of Cu. Fig. 4e and f lists the length distributions of both uncoated and Cu-coated CNTs in the powder mixtures. During milling process, the mechanical impact of stainless balls on CNTs and Al powders exerts a cutting effect, which shortened the CNTs [14, 39]. As a result, the average length of the attached CNTs (~830 nm for uncoated CNTs and ~800 nm for the Cu-coated CNTs, see Fig. 4e and f) is much shorter than that of raw CNTs (~3 µm). The final Al-(nano-Cu)-CNTs composite is obtained after sintering and hot-rolling. The electron backscatter diffraction (EBSD) inverse pole figures (IPF) maps in Fig. 5a-c confirm the

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single fcc structure of the matrix. The average grain sizes of the composite were counted based on the EBSD maps. The Al-CNTs composite (~1.9 µm, Fig. 5e) shows very similar grain size to the Al-(nano-Cu)-CNTs composite (~1.7 µm, Fig. 5f), of which grain sizes are slightly finer than the pure Al (~2.3 µm, Fig. 5d) due to the pinning effect of uniformly dispersed CNTs [40, 41].

Figure 4 (a) SEM micrograph of mixed Al flakes and uncoated CNTs. (b) SEM micrograph of mixed Al flakes and Cu-coated CNTs. (c) and (d) are high-magnification SEM micrographs for (a) and (b), respectively; inserted EDS results from the dash circle marked regions indicate the existence of Cu. (e) and (f) are length distributions of (c) uncoated CNTs and (d) Cu-coated CNTs in the mixing of Al powders and CNTs.

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Figure 5 EBSD inverse pole figure (IPF) maps of (a) pure Al, (b) Al-CNTs composite, and (c) Al-(nano-Cu)-CNTs composite. Corresponding grain size distribution of (d) pure Al, (e) AlCNTs composite, and (f) Al-(nano-Cu)-CNTs composite.

The bright field (BF) TEM image (Fig. 6) demonstrates that CNTs are uniformly dispersed in both the Al-CNTs composite and the Al-(nano-Cu)-CNTs composite. It is worth noting that no unwanted intermetallic compound is detected at the interfaces in current Al-(nano-Cu)-CNTs composite. Also, we can hardly detect any difference in interfacial structures between the AlCNTs composite and the Al-(nano-Cu)-CNTs composite based on regular TEM observations.

Figure 6 Bright field TEM images of (a) Al-CNTs composite and (b) Al-(nano-Cu)-CNTs composite.

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3.2 Interfacial structures To better unveil the interfacial structures of present composites, we further investigated the interfaces at atomic scales by means of the high-resolution scanning transmission electron microscopy (STEM). Fig. 7a-d shows typical Z-contrast images of the interfaces between Al and CNTs along <011> zone axis for the Al-CNTs (Fig. 7a and c) and Al-(nano-Cu)-CNTs (Fig. 7b and d) composites. Al displays dark contrast while Cu exhibits bright contrast. However, C is hard to be revealed in this Z-contrast imaging due to its low atomic number and the small fraction of C atoms in any atomic column along the electron beam direction. By comparing Fig. 7a with b, the Cu-rich transition region (~1-2 nm nanolayers) connecting Al and CNTs is clearly detected in the Al-(nano-Cu)-CNTs composite, whereas the Al-CNTs composite displays no such transition region. The EDS spectroscopy imaging in Fig. 8 further confirms the presence of Cu at the interface, in which <001> zone axis was reached for narrowing the transition region to achieve a better signal-to-noise ratio. Z-contrast imaging in Fig. 8a shows that atoms with light contrast are rich in the interfaces, and the corresponding EDS map in Fig. 8b indicates that these atoms are Cu.

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Figure 7 Z-contrast imaging of the interfaces in (a) Al-CNTs and (b) Al-(nano-Cu)-CNTs composites along <011> zone axis of Al. (c) and (d) are atomic resolution structures for interfaces corresponding to the marked regions in (a) and (b), respectively. (e) Schematic illustrations for interfacial structures of the composites. Cu-rich transition nanolayers are detected in (b), since Cu atoms exhibit brighter contrast due to their larger atomic number compared with Al atoms. The CNTs regions display slightly darker contrast. Yet the carbon atoms in CNTs are not directly visible since they merge in the Al matrix.

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The atomic resolution imaging at higher magnification distinctly demonstrates a transition of interfaces from severe loss of coherency (Fig. 7c) to ameliorative coherency (Fig. 7d). Such profound transition of interface structure is explained by the reduced lattice misfit strain resulting from the Cu-rich nanolayers. CNTs can be considered as a wrapped graphene with hexagonal structure (the space group of P6/mmm) [42]. Without considering the misfit strain, the most likely orientation relationship between face centered cubic (FCC) structured metals and hexagonal structures lies in <011>{111}FCC // <1010>{0001}Hexagonal, as detected in many other structural materials [43, 44]. However, a significant misfit strain (δ) does exist between <0001>CNTs and <111>FCC, which is calculated using [45] =2

−

(

+



)

(1)

where d represents the lattice spacing of the (h1k1l1) plane for FCC and (h2k2i2l2) plane for the CNTs, and m and n are the corresponding number of planes at the interface for the least misfit strain (m=6 and n =3 in this case). Given the lattice constant of Al (a=0.4046 nm) [46], Cu (a=0.3597 nm) [46] and CNTs (a=0.4772 nm, c=0.4129 nm) [42], the misfit strains are calculated as 12.33 % between <0001>CNTs and <111>Al, and 0.59 % between <0001>CNTs and <111>Cu.

Figure 8 (a) Z-contrast STEM imaging along <001>Al zone axis, showing the interfaces in the Al-(nano-Cu)-CNTs composite and (b) corresponding EDX mapping demonstrating Cu-rich transition regions in-between Al and CNTs.

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Based on these calculation results, the actual misfit strain between Al matrix and CNTs lies up to 12.33 % along the direction of <0001>CNTs// <111>FCC when CNTs are embedded within the Al matrix. This sorely suggests a deviation from this ideal orientation relationship, accounting for the experimental observation of incoherent interfaces between the Al matrix and CNTs (see Fig. 7c). Remarkably, Cu possesses smaller lattice constant than that of Al, reducing the misfit strain by an order of magnitude to 0.59 % in this specific direction. It serves as an ideal candidate to buffer the large interfacial misfit strain. During the processing procedure, the coated Cu atoms on the surfaces of CNTs diffused into Al matrix, forming the Cu-rich transition nanolayers, as shown in Fig. 7b. These buffering nanolayers reduce the large misfit between Al and CNTs, and consequently, substantially improve the coherency of interfaces, as shown schematically in Fig. 7e.

3.3 Mechanical performance Quasi-static tension tests (Fig. 9) were carried out to measure the mechanical performance of the Al-(nano-Cu)-CNTs composite, as well as the Al-CNTs composite and pure Al for comparison. By comparing with the Al-CNTs composite, the Al-(nano-Cu)-CNTs composite not only shows a raise of 54 MPa in yield strength and 96 MPa in tensile strength, but also exhibits an unexpected increase of 4.8 % in tensile elongation, indicating a breakthrough in the inverse strength-ductility relationship prevailed in structural materials [47, 48]. Overall, the Al-(nanoCu)-CNTs composite shows very high tensile ductility (elongation of 15.7 %), yet with a huge improvement of 170 MPa in yield strength in comparison to pure Al.

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Figure 9 Engineering stress-strain curves of pure Al, Al-CNTs, and Al-(nano-Cu)-CNTs composites. Both yield strength and tensile strength of the Al-(nano-Cu)-CNTs composite are higher than that of the Al-CNTs composite, along with better tensile elongation.

To figure out the strengthening effects of the strengthened interface, we have performed a quantitative analysis on the underlying strengthening mechanisms for yield strength of the composites. It can be described as following aspects concerning the increase in yield strength for Al-CNTs composites: (1) grain refinement [49], (2) Orowan looping strengthening of CNTs [50], and (3) load transfer from the matrix to CNTs [51]. As for the Al-(nano Cu)-CNTs composite, the solid solution strengthening of Cu atoms in Al matrix should also be included, because the substitutional diffusion of Cu into the Al matrix occurs during the composite fabrication process due to the similar atomic radius (Cu: 0.1278 nm, Al: 0.1434 nm) [52], resulting in the formation of Al-Cu solid solution (Cu-rich nanolayer) in the vicinity of CNTs. The strengthening effect from the grain refinement ( ∆σ GR ) is given by the Hall-Petch equation [39, 53]:

∆σ GR = K ( D −0.5 − D0−0.5 )

(2)

where D and D0 are the average grain sizes of composites and pure Al specimens, respectively. The value of Hall-Petch slope K is 0.04 MPa ∙ m0.5 for Al [54]. Based on the grain size information, the ∆σ GR for Al-CNTs and Al-(nano-Cu)-CNTs composites is 3 MPa and 4 MPa, respectively.

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The uniformly dispersed CNTs are able to inhibit movements of the nearby dislocations in Al matrix, and thus, improve the strength of the composite. Due to the high strength of CNTs (110 GPa), Orowan looping model is employed to estimate this part of yield strength contribution [55, 56]:

∆ σ Orowan =

φ 0.8MGb 1 . ln( ). 0.5 b λ −φ 2π (1 − v)

(3)

where M is the Taylor factor ( M = 3.06) [57], G is the shear modulus of the matrix (G=25.4 GPa) [58], b is the Burgers vector of the matrix (b=0.286 nm) [58], ν is Poisson's ratio (ν=0.33),

φ is the diameter of CNTs ( φ =10 nm), and λ is the inter-particle spacing ( λ =125 nm). λ was calculated based on the equation of λ = φ (0.5π / VCNTs ) 0.5 , since L is much larger than φ in this case, where L is the length of CNTs [42]. ∆σ Orowan for Al-CNTs and Al-(nano-Cu)-CNTs composites are both 107 MPa. With respect to Al-(nano Cu)-CNTs composite, the Cu atoms with different atomic sizes and shear modulus can cause a variation of strain fields in the nearby Al matrix of CNTs. Local strain fields are created that interact with dislocations and impede their motion, leading to an increase in the yield strength of the material. It has been generally accepted that the estimation of solidsolution strengthening (∆σss ) can be given by the Fleischer equation [59, 60]: 3

MGε ss2 c ∆σ ss = 700

(4) where ε ss = ε G′ − 3ε b

ε G′ =

εG 1 1+ ε G 2

εG =

1 dG G dc

εb =

1 da a dc

(5)

The meaning of M , G and b symbols in this equation are the same to those in eq. (3). ε ss is misfit parameter, a is the lattice parameter of the matrix (0.4046 nm for Al [46]), and c is the atomic concentration of the solute. By assuming that all the Cu (1.12 g) exists in the form of solid solution atoms throughout the Al matrix of Al-(nano Cu)-CNTs composite, therefore, c is 0.00364. The shear modulus and lattice parameter of Cu are 39 GPa and 0.3597 nm, respectively [61]. Combined with the shear modulus and lattice parameter of the matrix, ε G and ε b are 0.535 and -0.1064, respectively. Finally, ∆σ ss can be calculated as 4 MPa.

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Suppose the yield strength of composites could be regarded as the summation of σ M and strengthening contributions, the ∆σ LT can be estimated as σ YS − σ M − ∆σ GR − ∆σ Orowan for AlCNTs composite and σ YS − σ M − ∆σ GR − ∆σ Orowan − ∆σ ss for Al-(nano Cu)-CNTs composite. The values of ∆σ LT are 6 MPa and 55 MPa for Al-CNTs and Al-(nano-Cu)-CNTs composites, respectively. In addition, it is worth noting that the contribution of solid solution strengthening for Al-(nano-Cu)-CNTs composite should be less than 4 MPa, because the Cu solute atoms merely locate in the vicinity of CNTs rather than the entire Al matrix. So, the strength contribution from the enhanced load transfer for Al-(nano-Cu)-CNTs composite should be larger than 55 MPa. Fig. 10 summarizes the merits of each strengthening mechanism.

Figure 10 Theoretical calculations of yield strength contributions from each strengthening mechanism for pure Al, Al-CNTs, and Al-(nano-Cu)-CNTs composites. By comparison to the Al-CNTs composite, the improved yield strength for the Al-(nano-Cu)-CNTs composite mainly stems from the increase of the load transfer stress

In short, the improvement in yield strength from the grain refinement is small, i.e., 3 MPa and 4 MPa for the Al-CNTs and Al-(nano-Cu)-CNTs composites, respectively, since all three materials have the similar grain sizes (Fig. 5). As depicted in Fig. 10, the major strengthening effect of CNTs arises from their impediments to dislocation movements, which leads to an increase of 107 MPa in yield strength for both the Al-CNTs and Al-(nano-Cu)-CNTs composites, compared to pure Al.

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Regarding the discussion above, the improvement of yield strength mainly benefits from the higher load transfer, which increases from 6 MPa for the Al-CNTs composite to the 55 MPa for the Al-(nano-Cu)-CNTs composite. The raise of load transfer strength (49 MPa) in our composite is much higher than that of the composites processed by the interfacial reaction (the maximum increase of the yield strength is ~19 MPa) [23, 62]. This high load transfer is ascribed to the significantly improved interfacial bonding in the Al-(nano-Cu)-CNTs composite. In addition to the impressive increase of strength, the interfaces with better coherency enable higher extent of synergic deformation for Al matrix and CNTs upon loading, the so-called dynamic strain partitioning effect, which increases the critical damage nucleation strain [63, 64]. Coherent interfaces are highly desirable in pursuit of excellent mechanical performance for structural materials [47, 65-67]. Coherent nanoscale twin boundaries in ultrafine-grained Cu offer substantial strengthening effect and produce increase in ductility [47]. High densities of coherent nano-precipitates, for example, L12 particles in FCC matrix high entropy alloys [65] and B2 particles in maraging steels [67], enable ultrastrong structural materials while render excellent ductility. However, in metal matrix composites, most of the interfaces between reinforcements (e.g., fibers, whiskers, or particles) and matrix are incoherent [68], as also seen in our Al-CNTs composite. In this regard, we decorated Cu nanoparticles surrounding the CNTs and adapted a powder metallurgical routine to successfully sandwich the Cu-rich buffered layers in-between Al and CNTs. The buffered nanolayers reduced the interfacial misfit strain, leading to significant improvement in interfacial coherency and interfacial bonding. Therefore, a simultaneous increase in tensile strength and ductility is observed in the Al-(nano-Cu)-CNTs composite, as seen in Fig. 9. The difference in the mechanical properties of Al-CNTs and Al-(nano-Cu)-CNTs composites can also be justified by the fractural observation after tensile tests. As shown in Fig.11a and d, more dense and deeper dimples can be observed in the fracture surface of Al-(nano-Cu)-CNTs composite (Fig.11d), proving higher plastic deformation degree of Al matrix and better tensile ductility. In addition to different morphologies of Al matrix, the failure features of CNTs in the two types of composites are significantly different. As shown in Fig.11b, the tip of the pulled out CNTs is straight, implying the CNTs were directed extracted from the Al matrix during the tensile test (as displayed in the schematic image of Fig.11c). Unlike the CNTs morphologies in the fractural surface of Al-CNTs composite(Fig.11e), most of the pulled out CNTs in the

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fractural surface of Al-(nano Cu)-CNTs composite possess a sharp tip, denoting the CNTs appear to be broken (as displayed in the schematic image of Fig.11f) due to the efficient load transfer from matrix to CNTs. This is an evidence for the enhanced load transfer between Am matrix and CNTs by the interfacial nano-decoration.

Figure 11 The fracture surfaces and the deductive failure diagrams of (a, b, c) Al-CNTs composite and (d, e, f) Al-(nano Cu)-CNTs compsite

With a goal of tailoring the interfacial coherency, we succeed in producing an ultrahighperformance Al-(nano-Cu)-CNTs composite—the combination of tensile strength (391 MPa) and ductility (tensile elongation of 15.7 %) is better than its counterparts (see Fig. 12a). Furthermore, the mechanical performance of our composite outperforms many commercialized Al alloys, e.g., the tensile strength of our composite is higher than AA 6061 Al alloys (the highest strength is less than 350 MPa), and is comparable to high-strength Al alloys, such as AA2014 and AA 7075 Al alloys (strength of ~400-450 MPa for the underaged condition) [69]. More importantly, as demonstration of a new solution to enhance the interfaces in metal matrix composites, pure Al is used as the matrix in our study. Following our approach, Al alloys, instead of pure Al, can also be adopted as the matrix in the future. Thus, further improvement of the strength can be easily

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achieved, making the Al matrix composites promising candidates for high-performance lightweight structural materials.

Figure 12 Plot of (a) ultimate tensile strength (UTS) and (b) strengthening efficiency as a function of elongation to failure for CNTs reinforced pure Al matrix composites. The combination of UTS, strengthening efficiency and tensile elongation for our Al-(nano-Cu)-CNTs composite outperforms any other composites processed by methods of the high energy balling [70, 71, 73], the in-situ growth [7, 72, 74], and the medium energy ball milling with [23] and without [8, 12, 18, 55, 75-85] interfacial reaction.

The strengthening efficiency, a key index to evaluate the strength contribution of CNTs, was obtained by the ratio between the increase in tensile strength and the products of the tensile strength of matrix and the volume fraction of reinforcements [8]. The strengthening efficiency of CNTs (R) is expressed as [8]:

R=

σ c − σ TS VCNTsσ TS

(6)

21

where σ c and σ TS are the tensile strength of a given composite and Al matrix, respectively, and VCNTs is the volume fraction of CNTs. The strengthening efficiency is 106 for the Al-(nano-Cu)CNTs composite and 55 for the Al-CNTs composite. Fig. 12b compares the strengthening efficiency among various processing strategies as a function of tensile elongation. The high energy ball milling method has the advantage in acquiring high strengthening efficiency (115-125) in CNTs reinforced pure Al composites [70, 71], but the resulted highly distorted Al matrix leads to poor ductility (tensile elongation is typically less than 5 %) [71]. Insitu growth methods enable a good combination of strengthening efficiency (50-80) and tensile elongation (~17 %) [72]. However, the intrinsically incoherent interfaces between Al and CNTs limit the load transfer, and thus, retard further improvement of strengthening efficiency. Although the interfacial reactions are able to enhance interfacial bonding, they may severely damage the CNTs and produce unwanted compounds, which lower the strengthening efficiency to the range of 20-50 [23]. Unlike the previously proposed methods, our effort strikingly enhances the interfaces without damaging the integrity of CNTs. Having demonstrated that the present Al-(nano-Cu)-CNTs composite possesses Cu-transition region to improve the interfacial coherency, this composite shows high strengthening efficiency (105), while remains excellent tensile elongation (15.7 %).

4. Conclusions We present an interfacial nano-decoration approach to enhance the interfacial bonding of Al matrix composites. The interfaces with far better coherency are accomplished, which arise from the low interfacial misfit strain via forming Cu-rich transition nanolayers in-between Al and CNTs. The Al-(nano-Cu)-CNTs composite with an exceptional combination of tensile strength (391 MPa) and tensile elongation (15.7 %) was successfully fabricated by the processing combined with powder metallurgy method, which performs better than its peers, in terms of tensile strength, strengthening efficiency, and tensile elongation. The enhanced mechanical properties can be mostly attributed to the unique interface structure. Our strategy yields a strong benefit for the design of metal matrix composites with unprecedented mechanical properties to meet applications.

Acknowledgements

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We are grateful to Dr. Zhiming Li and Prof. Dierk Raabe at Max-Planck-Institut für Eisenforschung for their insightful comments. B. G., M. S. and Y. D. appreciate the financial support from National Natural Science Foundation of China (No. 51531009 and No. 51820105001), Key Laboratory of Advanced Materials of Yunnan Province (No. 738010094) and Postgraduate Innovation Program of Hunan Province (No. 150140012).

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Research highlights

Nano Cu particles were decorated onto the surface of CNTs to tailor the Al-CNTs interface. Cu-rich layer formed between Al and CNTs can increase the interfacial coherency. The strength and ductility of composite can be simultaneously increased by such tactic.

Declaration of interests ■ 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. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: