Al–Zn–Mg–Cu composites

Journal Pre-proof An investigation on particle weakening in T6-treated SiC/Al–Zn–Mg–Cu composites G.N. Ma, D. Wang, Z.Y. Liu, B.L. Xiao, Z.Y. Ma PII:

S1044-5803(19)31641-9

DOI:

https://doi.org/10.1016/j.matchar.2019.109966

Reference:

MTL 109966

To appear in:

Materials Characterization

Received Date: 18 June 2019 Revised Date:

13 September 2019

Accepted Date: 13 October 2019

Please cite this article as: G.N. Ma, D. Wang, Z.Y. Liu, B.L. Xiao, Z.Y. Ma, An investigation on particle weakening in T6-treated SiC/Al–Zn–Mg–Cu composites, Materials Characterization (2019), doi: https:// doi.org/10.1016/j.matchar.2019.109966. 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 Inc.

An investigation on particle weakening in T6-treated SiC/Al-Zn-Mg-Cu composites G. N. Maa,b, D. Wanga,*, Z. Y. Liua, B. L. Xiaoa,*, Z. Y. Maa a

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b

School of Materials Science and Engineering, University of Science and Technology of China, 72 Wenhua Road, Shenyang 110016, China

Abstract This study found weakening effect of SiC particles on strength of a T6-treated SiC/Al-Zn-Mg-Cu composite. The mechanisms for the weakening were investigated in detail for the first time. The 15 vol.% SiC/7085Al composite and 7085Al alloy were fabricated by powder metallurgy (PM) method and subjected to microstructural examination and mechanical properties evaluation. It was indicated that the as-extruded composite had 92 MPa enhancement in the ultimate tensile strength (UTS) compared to the as-extruded unreinforced alloy. However, After T6 treatment (solution treatment and artificial ageing), the UTS of the composite was 34 MPa lower than that of the unreinforced alloy. Microstructural examination disclosed that the content of η′-MgZn2 was reduced in the T6-treated composite due to interface reaction between Mg and oxide on the surface of SiC particles. Meanwhile, compared to the unreinforced alloy, the composite had coarser precipitates with relative lower density under the T6 condition. Furthermore, it was found that SiC particles accelerated ageing process, leading to formation of coarse precipitates and precipitate free zone in the

*

Corresponding author. Tel./fax: +86-24-23971749, E-mail address: [email protected] (D. Wang); Tel.: +86-24-83978630, Fax: +86-24-23971749, E-mail address:, [email protected] (B.L. Xiao). 1

vicinity of SiC-Al interface. According to the theoretical calculation, the weakened precipitation strengthening value was not compensated by mismatch dislocation strengthening and load transfer strengthening. By adding additional Mg (1.0 wt.%) into the composite, the precipitation strengthening effect could be basically recovered, resulting in significantly enhanced strength compared to the 7085Al alloy. Keywords: metal matrix composites, mechanical properties, interfacial reaction, ageing, strengthening mechanisms 1. Introduction Particle reinforced metal matrix composites (PRAMCs) have been widely used in various industrial fields due to their high specific strength and specific modulus, low coefficient of thermal expansion, etc [1-3]. Significant enhancement of strength exerted by particles such as silicon carbide (SiC) has been verified in 2xxx, 5xxx and 6xxx Al alloys [4-6]. In general, stronger matrix alloys tend to produce stronger composites, therefore, high strength Al alloys were usually chosen as the matrices for PRAMCs as structural materials. Most of Al-Zn-Mg-Cu alloys have higher strength than 2xxx, 5xxx and 6xxx Al alloys, such as 7075Al, 7085Al and 7055Al alloys. It was expected that ultra-high strength PRAMCs were developed based on these Al-Zn-Mg-Cu alloys. Unfortunately, the strengthening effect of SiC particles in 7xxx Al alloys was far from in 2xxx Al alloys [7-10]. Even SiC particles exhibited weakening effect in SiC/Al-Zn-Mg-(Cu) composites, i.e. the composites had lower strength than the unreinforced alloys [11-13]. Lee et al [7] reported that the mismatch between reinforcement and matrix leads to a large stress 2

concentration near the reinforcement, and the matrix in that region fails prematurely, resulting in a decrease of yield strength. Min et al [10] supposed that the weak matrix and interface debonding could lead to insignificant strengthening effect of SiC particles. Considering these mechanisms also exert in PRAMC based on other aluminium alloys, etc., the weakening mechanisms of SiC particles in Al-Zn-Mg-Cu alloys might be unknown. Hong et al [14] investigated the effect of SiC/Al interfacial reaction on ageing behavior in detail, however, the reason for the weakening effect of SiC particles on the mechanical properties was not explained. So far, the relevant weakening mechanisms have been not well understood. Compared to 2xxx and 6xxx Al alloys, the precipitation strengthening effect of Al-Zn-Mg-Cu alloys was more significant. Naturally, the main precipitation phase (MgZn2) in Al-Zn-Mg-Cu alloys has significant effect on mechanical properties of their composites. Mg is extremely active and able to react with the oxide and other impurities introduced by SiC particles and raw metal powders to form Mg2Si, MgO and MgAl2O4 during fabrication process [15,16]. This could reduce the content of MgZn2. Meanwhile, SiC particles could also change precipitating behaviors of matrix alloys. Hong et al [14] reported that SiC particles decelerated ageing process of SiC/Al-7.0Zn-2.0Mg-2.0Cu composite. They considered that the SiC-Al interface reaction depleted solute atoms, which reduced driving force for heterogeneous nucleation of precipitates. On the contrary, Ma et al. [17] found an accelerated ageing process in SiC/7075Al composites. They explained that SiC particles provided more preferential nucleation sites for precipitation and reduced the activation energy for diffusion of solid solution atoms. Moreover, the dislocations and lattice distortions introduced by SiC 3

particles could accelerate abnormal growth of partial precipitation phases. Clearly, the factors affecting the mechanical behavior of SiC/Al-Zn-Mg-Cu composites are quite complicated and no consistent conclusions have been reached so far. In-depth studies are highly needed for controlling of strength of SiC/Al-Zn-Mg-Cu composites. In this study, 15 vol.% SiC/7085Al composite was fabricated by powder metallurgy (PM) technique. Microstructure and mechanical properties of both the composite and unreinforced alloy were compared. The aim is to (a) understand the effect of SiC particles on the microstructure of SiC/Al-Zn-Mg-Cu composites, (b) elucidate the negative factors influencing strength of the composites, and (c) provide the basis for enhancing mechanical properties of SiC/Al-Zn-Mg-Cu composites. 2. Experimental 2.1. Fabrication of SiC/7085Al composite 15 vol.% SiC/7085Al composite was fabricated using the PM technique. The 7085Al alloy had a nominal composition of Al-7.5Zn-1.8Mg-1.7Cu (wt.%). Considering Mg consuming due to interface reaction, Al-7.5Zn-2.8Mg-1.7Cu (7085Al with surplus 1.0 wt.% Mg) was also adopted to be matrix alloy. The mean sizes of alloy powder (99.9 pct. purity, Changsha Tianjiu Metal Material Co., Ltd.) and SiC particles (99.5 pct. purity, Baige Group Co., Ltd.) were 13 and 7 µm, respectively. It can be seen from Fig. 1(a) that the SiC particles had a polygonal morphology with sharp edge and low aspect ratio. The main impurities in the SiC particles were free Si and SiO2 as shown in Fig. 1(b).

4

Fig. 1 (a) Morphology and (b) XRD analysis of the as-received SiC particles. The SiC particles were mixed with Al powder for 7 h using a bi-axis rotary mixer with a rotation speed of 50 rpm. The mixed powders were degassed in a vacuum furnace under 10-2 Pa, and sintered at 560 ºC for 2 h, then hot pressed into billets in a die with a height of 300 mm and a diameter of 120 mm under a pressure of 70 MPa. For comparison, the unreinforced 7085Al alloy and SiC/Al-7.5Zn-2.8Mg-1.7Cu (SiC/7085Al-1.0Mg) composite were also fabricated using the same method. The nominal compositions of three materials are listed in Table 1. The billets were heated at 420 ºC for 1 h and were extruded into bars with an extrusion ratio of 17:1. Part of the extrusion bars were T4-treated (solution treated at 470 ºC for 2 h, water quenched) and T6-treated (aged at 120 ºC for 24 h after T4-treated). Table 1. Nominal compositions of the unreinforced alloy and the matrices of the composites. Materials

wt.%

vol.%

Zn

Mg

Cu

Al

SiC

7085Al

7.5

1.8

1.7

Bal.

N/A

SiC/7085Al

7.5

1.8

1.7

Bal.

15.0

SiC/7085Al-1.0Mg

7.5

2.8

1.7

Bal.

15.0

5

2.2. Mechanical property tests Tensile specimens with a gauge length of 20 mm, a width of 3.8 mm and a thickness of 2.3 mm were machined from the extruded bars with the tensile axis parallel to the extrusion direction. Tensile test were conducted at a strain rate of 1×10-3 s-1 and room temperature using an Instron 5848 tester. At least 5 tensile specimens were tested for each the material. In order to evaluate the effect of SiC particles on solid-solution and precipitation strengthening, the tensile specimens under as-extruded condition and as-solution treated (T4-treated) condition were also tested for each material in the same way. 2.3. Characterization of microstructure and microchemistry The specimens for microstructural examinations were cut along the extrusion direction. The specimens were ground with 2000 grit abrasive paper, mechanically polished, corroded by Keller′s etchant and then examined by optical microscopy (OM, Leica). According to Archimedes principle, the densities of the unreinforced alloy and composites were measured using solid density meter (KW-120E). The theoretical densities were calculated by mixing rule. Phases were identified using an X-ray diffractometer (XRD, D/max 2400). The interface and precipitation phase were examined by transmission electron microscopy (TEM, TECNAI G2 F20). The thin foils for TEM were mechanically polished and ion-milled. Chemical compositions of the composites and unreinforced alloy were analyzed by wave dispersive spectrometer (WDS). The element distribution in the composites was examined by electron probe micro analysis (EPMA-8530F). The fracture surfaces of the SiC/7085Al composite and 7085Al alloy under the T6 condition were examined by scanning electron 6

microscopy (SEM, quanta 600). 3. Results 3.1. Microstructure of SiC/7085Al composite Table 2 shows the densities of the as-extruded composites and 7085Al alloy. All the composites and unreinforced alloy were densified, and no porosities were observed. Fig. 2 shows that the mean grain sizes of the 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composite were 3.4, 3.1 and 3.8 µm, respectively. For the materials prepared by the PM technique, the difference in the mean grain size between the alloy and composites is usually small. SiC particles were distributed randomly with their major axis being parallel to the extrusion direction (Fig. 2(b, c)). Particles could rotate with plastic flowing of matrix alloy during thermo-mechanical processing, and cusps of SiC particles perpendicular to the extrusion direction tended to be broken up [18,19]. As a result, the SiC particles in the composites exhibited an orientation distribution and finer size compared to the as-received particles (Fig. 1(a)). Table 2. Density of the as-extruded composites and unreinforced alloy. Measured density

Theoretical density

(g.cm-3)

(g.cm-3)

7085Al

2.833

2.838

99.80

SiC/7085Al

2.885

2.896

99.60

SiC/7085Al-1.0Mg

2.872

2.879

99.76

Materials

7

Relative density (%)

Fig. 2 OM images and grain size statistic plots of (a) the 7085Al alloy, (b) the SiC/7085Al composite and (c) the SiC/7085Al-1.0Mg composite under T6 condition. 3.2. Mechanical properties The tensile curves of the T6-treated 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites are shown in Fig. 3. The work-hardening process of the composites was more obvious than that of the unreinforced alloy. Unlike the other PRAMCs based on 2xxx or 6xxxAl alloys, the SiC/7085Al composite had significantly lower ultimate tensile strength (UTS) and yield strength (YS) than the unreinforced alloy.

Fig. 3 Stress-strain curves of the composites and unreinforced alloy. Fig. 4(a) shows that particle fracture and matrix alloy fracture occurred in the SiC/7085Al composite, and no interfacial debonding was observed, mainly due to good interface bonding. However, compared to the fracture morphology of the 7085Al alloy (Fig. 4(b)), less and shallower dimples in the SiC/7085Al composite implied low ductility of the 8

composite. Based on the density and fractography analysis, it was not metallurgical quality that led to the reduced strength of the SiC/7085Al composite.

Fig. 4 Fractographs of (a) the SiC/7085Al composite and (b) the 7085Al alloy under T6 condition. The tensile properties of the unreinforced alloy and composites under different conditions are presented in Table 3. The UTS of the as-extruded SiC/7085Al composite was 92 MPa higher than that of the as-extruded 7085Al alloy, while the T4-treated SiC/7085Al composite had almost same UTS as the T4-treated 7085Al alloy. As shown in Table 3, after solution treatment, the UTS of the 7085Al alloy was improved by 128 MPa, but the UTS of the SiC/7085Al composite was improved by only 36 MPa. Moreover, after T6 treatment, the UTS of the 7085Al alloy was improved by 182 MPa, while only an UTS increment of 148 MPa was observed for the SiC/7085Al composite. Although the YS of the 7085Al alloy and SiC/7085Al composite both were enhanced, the YS increment of the SiC/7085Al composite (257 MPa) was significantly smaller than that of the 7085Al alloy (316 MPa). Finally, the UTS of the T6-treated SiC/7085Al composite was 34 MPa lower than that of the T6-treated 7085Al alloy. The results indicate that strengthening of 9

SiC particles in the as-extruded, T4 and T6 treated composites were significant, weak and negative, respectively. This phenomenon was never expected and would be discussed in detail in the latter part of the paper. Table 3. Tensile properties of the composites and unreinforced alloy under different conditions. Materials

Condition

UTS (MPa)

YS (MPa)

Failure strain (%)

7085Al

As-extruded

277±11

153±6

15.2±0.3

T4

405±1

242±4

12.7±0.3

T6

587±7

558±8

11.1±0.9

As-extruded

369±2

204±3

8.0±0

T4

405±5

229±4

8.6±0.6

T6

553±7

486±9

3.5±0.5

As-extruded

357±1

205±3

5.7±0.7

T4

455±2

288±4

5.5±0.5

T6

663±4

617±7

2.3±0.6

SiC/7085Al

SiC/7085Al-1.0Mg

Table 3 also shows that the as-extruded SiC/7085Al-1.0Mg composite had similar UTS and YS to the as-extruded SiC/7085Al composite. That is, the addition of extra 1.0 wt.% Mg did not increase the strength of the as-extruded SiC/7085Al composite. The UTS increments after heat treatment in the SiC/7085Al-1.0Mg composite were larger than that in the SiC/7085Al composite. The improvements of the UTS of the SiC/7085Al-1.0Mg composite after T4 and T6 treatment were 98 MPa and 208 MPa, respectively, which was close to that of the 7085Al alloy.

3.3. Chemical reaction products in the composites Chemical compositions disclosed by WDS are shown in Table 4. Mg content in the region away from the SiC-Al interface of the SiC/7085Al and SiC/7085Al-1.0Mg composites 10

was 0.7 wt.% and 1.0 wt.% lower than the nominal content (Table 1), respectively. The SiC/7085Al-1.0Mg composite had similar matrix composition to the 7085Al alloy. Fig. 5 shows XRD patterns of the as-extruded 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites. The peaks of MgZn2 were detected in the three materials. Moreover, the peaks of Mg2Si were also identified in the SiC/7085Al and SiC/7085Al-1.0Mg composites. Table 4. Chemical compositions in the matrix region of the T6-treated composites and unreinforced alloy determined by EPMA (wt.%). Materials

Al

Mg

Zn

Cu

7085Al

88.8±0.41

1.9±0.09

7.6±0.08

1.7±0.09

SiC/7085Al

89.6±0.30

1.1±0.26

7.6±0.23

1.7±0.10

SiC/7085Al-1.0Mg

88.7±0.34

1.8±0.12

7.9±0.11

1.6±0.15

Fig. 5 XRD patterns of the as-extruded composites and unreinforced alloy. The rod-like Mg2Si with sizes of 300-800 nm was found in the SiC/7085Al composite (Fig. 6(a)), which was in accordance with the results shown in Fig. 5. Fig. 6(b) shows that a small amount of interface products with sizes of 30-50 nm were discontinuously distributed at the SiC-Al interface. The interface products had been generally considered as MgO or MgAl2O4 [14,20], which needed to be further verified using high-resolution transmission 11

electron microscopy (HRTEM). Fig. 6(c-g) indicates that the interface products could be determined to be face centered cubic MgO (Fm-3m space group, a=b=c=0.421 nm), which had a crystallographic orientation relationship with the SiC particle, (011)SiC∥(011)MgO. Fast Fourier transform (FFT) recorded from the red dotted rectangles shows diffraction spots of the —

SiC along the [011] zone axis (Fig. 6(e)), and the inter-planar spacing on the plane of (200)SiC —

and (111)SiC was measured as 0.221 and 0.256 nm, respectively (Fig. 6(c)). Fig. 6(f) shows the clear crystal lattices with the measured inter-planar spacing of 0.247 and 0.210 nm, which —

—

correspond to the plane of (111)MgO and (200)MgO, respectively.

Fig. 6 Bright-field TEM images of (a) Mg2Si phase and (b) SiC-Al interface in the T6-treated SiC/7085Al composite; (d) HRTEM image of interface product; (c) enlarged view and (e) FFT image corresponded to red dotted rectangle in (d) image, (f) magnified view and (g) FFT 12

image corresponded to yellow dotted rectangle in (d) image. 3.4. Element distribution of SiC/7085Al composite Fig. 7(a-d) shows the SEM image and elemental distribution maps of the as-extruded SiC/7085Al composite. The backscattered image indicates that many MgZn2 particles with sizes of 2-4 µm were randomly distributed in the Al matrix. Most of MgZn2 phases had formed when alloy powders were prepared. Actually, during sintering process, intermetallics could also fully form [21]. Their locations were in consistence with Zn, Mg rich zones (Fig. 7(a)). Cu atoms also were enriched in these zones (Fig. 7(d)) because Cu could dissolve in MgZn2 phase to replace Zn atom position [22]. After T6 treatment, the MgZn2 particles were refined and hardly observed (Fig. 7(e)). Most of the Zn and Mg rich intermetallics were homogenously distributed in the matrix region, except for a small amount of Mg aggregation around SiC particles (Fig. 7(f, g)). Furthermore, a few of Al7Cu2Fe identified by WDS existed pointed by the arrows because their higher solution temperature in Al matrix. 3.5. Precipitated phase The bright field TEM images and corresponding selected area electron diffraction (SAED) along [110]Al direction for the three materials under T6 condition are shown in Fig. 8(a, b, c). The needle-like precipitates were identified to be the η′-phase, which had obviously different densities and sizes in the three materials. Compared to the 7085Al alloy, the precipitates in the SiC/7085Al composite had lower density and larger size. In the SiC/7085Al-1.0Mg composite, the density of the precipitates was similar to that in the 7085Al 13

alloy, whereas the size of the precipitates was close to that in the SiC/7085Al composite. Fig. 8(d, e, f) shows that the precipitates in the 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites had a mean length of 6.2, 10.3 and 8.6 nm, respectively.

14

Fig. 7 SEM images and elemental maps of (a, b, c, d) the as-extruded SiC/7085Al composite and (e, f, g, h) the T6-treated SiC/7085Al composite: (a, e) backscattered images, (b, f) Zn 15

maps, (c, g) Mg maps and (d, h) Cu maps.

Fig. 8 Bright-field TEM images and precipitate size statistic plots of (a, d) the 7085Al alloy, (b, e) the SiC/7085Al composite and (c, f) the SiC/7085Al-1.0Mg composite under T6 condition taken along [110] Al zone axis. Near the SiC-Al interfaces, the precipitates in the T6-treated SiC/7085Al composite were heterogeneously distributed (Fig. 9). In regions I and II marked by red dash line, coarse precipitates were found, the length of the longest precipitates was about 38.6 nm. Meanwhile, the densities of the precipitates in these regions were significantly lower. Moreover, there existed a discernible precipitate free zone (PFZ) with a width of ~40 nm adjacent to the SiC-Al interface (Fig. 10).

16

Fig. 9 TEM image showing the precipitates at SiC-Al interface of the T6-treated SiC/7085Al composite.

Fig. 10 PFZ adjacent to SiC-Al interface of the T6-treated SiC/7085Al composite. 4. Discussion 4.1. Effect of SiC particles on microstructures In the present composites (Fig. 5), Mg2Si was unexpectedly found because Si and SiO2 in the SiC powders (Fig. 1(b)) could initiate the following chemical reactions [23,24]: 2Mg+Si — Mg2Si 17

(1)

2Mg+ SiO2 — 2MgO+Si Mg+ 2SiO2 +2Al — MgAl2O4+2Si

(2) (3)

HRTEM images shown in Fig. 6(d) confirmed that the interface reaction product was MgO. The formation of Mg2Si and MgO consumed Mg, thereby reducing the MgZn2 content. Meanwhile, Fig. 7(g) shows that Mg distribution was nonuniform. Segregation of Mg at the SiC-Al interfaces furtherly decreased the Mg content in the matrix far from the interfaces. Aggregation of Mg near the interfaces between Al and reinforcements was commonly observed in SiC/Al-Cu-Mg composites [21,25] and CNT/Al-Mg-Si composites [26]. At the interfaces, excess vacancies were highly beneficial to Mg diffusion [25]. Meanwhile, interface reactions could induce Mg diffusion from the matrix to the interfaces. As shown in Fig. 8(a, b), compared to the 7085Al alloy, the SiC/7085Al composite had coarser and lower density precipitates. Similarly, in regions I and II shown in Fig. 9, the precipitates were coarser and sparser. The results suggested that ageing was accelerated by the SiC particles. Lattice defects introduced by SiC particles would act as preferential nucleation sites for heterogeneous nucleation. Owing to the coefficient of thermal expansion mismatch between SiC and Al, dislocation density in the matrix near the SiC-Al interface was increased, and then dislocations accelerated atom diffusion and precipitate coarsening. Furthermore, because Mg was consumed to form Mg2Si and MgO, Zn/Mg ratio in the matrix was increased, excess Zn could facilitate η′-phase growth. After ageing treatment, the PFZ was noted in the vicinity of the SiC-Al interfaces (Fig. 10). The PFZs usually were formed in the vicinity of grain boundaries due to vacancy or 18

solute depletion mechanism [27,28]. It could be postulated that a similar mechanism would act in the case of the SiC-Al interfaces. The coarser precipitates preferentially formed at the interfaces produced poor solute atom zones near the interfaces, leading to formation of PFZ. 4.2. Effect of SiC particles on strengthening mechanisms A large number of investigations [29-31] revealed that the strengthening mechanisms of metal matrix composites include grain boundary strengthening, solid solution strengthening, precipitation strengthening, dislocation strengthening and load transfer strengthening. In order to understand the weakening phenomenon in the T6-treated SiC/7085Al composite, the varied strengthening mechanisms were evaluated. Quantitative estimates of the contribution of the major mechanisms in the 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites were described using a series of equations. The physical meaning and values of the symbols in the equations are summarized in Table 5. 4.2.1 Grain boundary strengthening As shown in Fig. 2, the present composites and unreinforced alloy exhibited quite similar grain size. Hence, grain boundary strengthening between the unreinforced alloy and the composites were similar. The increase in the yield strength (YS) due to the grain boundary strengthening ∆σgb could be given by a Hall-Petch-type relationship [32]:

∆σ gb =k y d −1/2

(4)

where ky is Hall-Petch coefficient and listed in Table 5, d is the mean grain size. From Eq. (4), ∆σgb of the T6-treated 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites could be calculated to be about 65, 68 and 62 MPa, respectively. 19

Table 5. Physical meaning and values of different symbols used in strengthening mechanism calculations. Symbol

Meaning

Values

Reference

M

Mean orientation factor

3.06

[32]

G

Shear modulus

26.9

[32]

b

Burgers vector

0.286 nm

[32]

ν

Poisson ratio

0.33

[32]

ky

Hall-Petch coefficient

0.12 MPa m-1/2

[32]

20, 38, 25 nm

Measured values

Mean precipitate spacing (7085Al alloy, λp

SiC/7085Al and SiC/7085Al-1.0Mg composites)

Vp

Volume fraction of SiC particle

0.15

Present work

β

Taylor’s coefficient

1.25

[36]

19.5 10-6K-1

[36]

450 K

Present work

Difference between coefficient of ∆CTE

thermal expansion of matrix alloy and SiC particle

∆T

Difference between processing and test temperature

dp

Average size of SiC particle

7.0 µm

Present work

S

The aspect ratio of particulates

2.5

Present work

4.2.2 Solid solution strengthening As solute atoms, Zn, Mg and Cu have different sizes and shear modulus from Al, which can cause a variation of strain fields, leading to an increase in the YS of Al-Zn-Mg-Cu alloys [32]. The strength increase induced by solid solution strengthening ∆σss could be described by the Fleischer equation [33]: ∆ σ ss = M G b ε 3 / 2

c

(5)

where M, G and b are listed in Table 5. From Eq. (5), the value of ∆σss depends on the concentration of the solute atoms c and the lattice strain ε. For Al-Zn-Mg-Cu alloys, the 20

theoretical contributions of Zn, Mg and Cu atoms per mass fraction to the YS were calculated to be 2.9, 18.6 and 13.8 MPa wt.%-1 [32]. The actual compositions in the matrix of the composites and unreinforced alloy are showed in Table 4. After quenched, assuming that all the alloyed elements are solute atoms and the contribution of the different solute atoms are additive, the solution strengthening in the 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites accounts for a strength increment of 81, 67 and 78 MPa, respectively. The calculated results are the upper bound of the solution strengthening in the three materials. Under T6-treated condition, η′-MgZn2 phases formed and part of Mg segregated to the SiC-Al interfaces. Thus, the actual difference in the solution strengthening between the T6-treated 7085Al alloy and SiC/7085Al composite should be less than 14 MPa. 4.2.3 Precipitation strengthening Orowan dislocation bypassing is the operative mechanism for GP zones and η′-phase [32]. For the T6-treated 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites, the main second phases are GP zones and η′-phase. Thus, the YS increment, ∆σor, could be calculated by the following equation [34,35]:

∆σ or =

0.4GMb ln(1.63r / b) λp π 1 −ν

(6)

where ν and λp are listed in Table 5; r is the mean radius of precipitates. For simplified calculation, the precipitates could be supposed to be spherical, r = l/π, where l is the mean length of the needle-like precipitates, 6.3, 10.2 and 8.6 nm, corresponding to the T6-treated 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites, respectively. As shown in Fig. 8, the spacing of the precipitates in the T6-treated SiC/7085Al 21

composite was much larger than that in the unreinforced alloy. Therefore, ∆σor of the T6-treated SiC/7085Al composite might be lower. From Eq. (6), ∆σor of the T6-treated 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites could be estimated to be about 446, 283 and 402 MPa, respectively. For the 7085Al alloy, SiC/7085Al and SiC/7085Al-1.0Mg composites, after artificial ageing treatment, the YS were enhanced 316, 257 and 329 MPa (Table 3), respectively. Assuming that the increments of YS were experimental values of precipitation strengthening. It is obvious that the experimental values were less than the estimated values. The experimental values obtained by this way was smaller because the loss of the solution strengthening under T6-treated condition was ignored. Moreover, during measuring λp, the sample thickness was not considered. The precipitates on different planes could be projected onto a TEM image. Thus, the λp measurement values could be smaller, resulting in the larger estimated values of precipitation strengthening. Anyhow, the estimated and experimental results all indicated that precipitation strengthening effect was restricted due to the addition of SiC particles, which should be the main reason why the YS of the SiC/7085Al composite was lower than that of the 7085Al alloy under T6-treated condition. 4.2.4 Mismatch dislocation strengthening To analyze the difference in dislocation strengthening mechanism of the alloy and composites, the increased density of geometrically necessary dislocations (GNDs) resulting from the coefficient of thermal expansion (CTE) mismatch due to quenching was only considered. The contribution by the GND strengthening ∆σgnd could be estimated by 22

following equation [36-38]: ∆σ

gnd

= β Gb

12 2Vp ∆ CTE ∆ T (1−Vp )bdp

(7)

where β, ∆CTE, dp and ∆T are listed in Table 5. GNDs usually gather inside the punched zone around SiC particles. When the volume fraction of SiC particles is higher than 25%, the dislocation punched zone of a particle will touch that of a neighboring particle and Eq. (7) may be inapplicable [38]. In this experiment, the volume fraction of SiC particles is 15%, thus Eq. (7) is used. From Eq. (7), it can be seen that the mismatch dislocation strengthening is related to Vp and dp and unrelated to composition of the matrix alloy. Thus, the mismatch dislocation strengthening values of the SiC/7085Al and SiC/7085Al-1.0Mg composites should be same. Using the values of Table 5 in Eq. (7), the mismatch dislocation strengthening values in the SiC/7085Al and SiC/7085Al-1.0Mg composites were estimated to be 34 MPa. 4.2.5 Load transfer strengthening Compared to that in the unreinforced alloy, the load transfer strengthening is a unique strengthening mechanism in the composites. In the SiC/7085Al and SiC/7085Al-1.0Mg composites, assuming that all the SiC-Al interfaces were well bonded, the strengthening due to the load transfer effect ∆σl-t could be estimated by the following equations [36]: ∆σ l-t =S/2Vpσ p-m

σ p-m =σ 0 + ∆σ gb + ∆σ or + ∆σ gnd

(8) (9)

where Vp and S is listed in Table 5; σp-m is the YS of the plastic matrix zone near the SiC particles; σ0 is the YS of the pure Al in single crystal, σ0 =17 MPa. From Eq. (8), ∆σl-t of the 23

T6-treated SiC/7085Al and SiC/7085Al-1.0Mg composites could be calculated to be about 75 and 97 MPa, respectively. This result indicated that weakened precipitation strengthening reduced load transfer strengthening. According to the modified shear lag model, YS of the materials, σc, could be calculated by the following equations [39]:

σ c =σ PVP +σ p-mVp-m +σ mVm=σ p-m [VP( S+2 )/ 2] + σ p-m ( R3 − 1)VP + σ m (1 − R3VP ) (10) σ m=σ 0 + ∆σ gb + ∆σ or

(11)

where σp is the stress that the SiC particles transfers, σm is the YS of the elastic matrix zone near the SiC particles [39], R is the ratio of the diameter of the plastic zone to that of the SiC particles, R=1.29 [40]. For the T6-treated 7085Al alloy, σc=σm. The estimated YS of the T6-treated composites and unreinforced alloy are listed in Table 6. In the T6-treated composites and unreinforced alloy, ∆σss was minor and not considered for estimating the YS of the T6-treated composites and unreinforced alloy. According to Eq. (8 and 9), ∆σl-t might be underestimated. Thus, it is reasonable that the estimated values were slightly lower than the measured values. Consistent with the measured results, the estimated YS of the T6-treated SiC/7085Al composite was also lower than that of the T6-treated 7085Al alloy. By comparing these above strengthening mechanisms listed by Table 6, it is obvious that the precipitation strengthening was the main strengthening mechanism in the T6-treated 7085Al alloy and SiC/7085Al composite. SiC particles had both positive and negative effects on the YS of SiC/7085Al composite. The weakened precipitation strengthening (163 MPa) counteracted the load transfer strengthening (75 MPa) and mismatch dislocation strengthening (34 MPa) 24

effect of SiC particles, resulting in lower strength of the T6-treated SiC/7085Al composite. When 1.0 wt.% additional Mg was added into the SiC/7085Al composite, the precipitation strengthening was enhanced, therefore the strengthening effect of SiC particles was revealed. Choosing Al-Zn-Mg-Cu alloys with high Mg content as matrix alloy is an effective method to fabricated high strength composites. In addition, other methods (e.g., SiC particles surface treatment) to avoid weakening precipitation strengthening are worthy to be researched in the future. Table 6. Estimated strength increment for different strengthening mechanisms in the T6-treated composites and unreinforced alloy. 7085Al

SiC/7085Al SiC/7085Al-1.0Mg

∆σgb (MPa)

65

68

62

∆σss (MPa)

<<81

<<67

<<78

∆σor (MPa)

446

283

402

∆σl-t (MPa)

/

75

97

∆σgnd (MPa)

/

34

34

σc (MPa)

528

460

594

σys (MPa)

558

486

617

σc: yield strength estimated by Eq. (10). σys: yield strength measured by tensile test. 5. Conclusions In the present work, the T6-treated SiC/7085Al composite exhibited lower tensile strength compared to the unreinforced alloy. Based on the experimental investigations and calculation of various strengthening mechanisms, the following conclusions can be made: (1) Mg reacted with the impurities (SiO2 and Si) from SiC powders to form MgO and Mg2Si, resulting in significantly reduced capability of the precipitation strengthening of the T6-treated SiC/7085Al composite. 25

(2) SiC particles accelerated the ageing process, the precipitates in the SiC/7085Al composite was larger and sparser than those in the 7085Al alloy. Meanwhile, the precipitates were nonuniform, and partial coarsening and PFZs existed in the vicinity of the SiC-Al interfaces. (3) After ageing treatment, the YS increment of the SiC/7085Al composite was significantly smaller than that of the 7085Al alloy. The weak precipitation strengthening was the main reason for the lower YS increment of the T6-treated SiC/7085Al composite. (4) By adding additional Mg (1.0 wt.%) into the SiC/7085Al composite, the negative effect of SiC particles was basically eliminated, and mechanical properties of the SiC/7085Al composite were significantly improved. Acknowledgments The authors gratefully acknowledge support of the National Key R & D Program of China (No. 2017YFB0703104) and the National Natural Science Foundation of China (grant Nos. 51771193, U1508216).

References [1] J. Song, Q. Guo, Q. Ouyang, Y. Su, J. Zhang, E.J. Lavernia, J.M. Schoenung, D. Zhang, Influence of interfaces on the mechanical behavior of SiC particulate-reinforced Al-Zn-Mg-Cu

composites,

Mater.

Sci.

Eng.

A

644

(2015)

79-84.

https://doi.org/10.1016/j.msea.2015.07.050 [2] Z.Y. Ma, S.C. Tjong, L. Gen, In-situ Ti-TiB metal-matrix composite prepared by a reactive

pressing

process,

Scr. 26

Mater.

34

(2000)

367-373.

https://doi.org/10.1016/S1359-6462(99)00354-1 [3] J.F. Zhang, X.X. Zhang, Q.Z. Wang, B.L. Xiao, Z.Y. Ma, Simulations of deformation and damage processes of SiCp/Al composites during tension, J. Mater. Sci. Technol. 34 (2018) 627-634. https://doi.org/10.1016/j.jmst.2017.09.005 [4] K. Amouri, S. Kazemi, A. Momeni, M. Kazazi, Microstructure and mechanical properties of Al-nano/micro SiC composites produced by stir casting technique, Mater. Sci. Eng. A 674 (2016) 569-578. https://doi.org/10.1016/j.msea.2016.08.027 [5] A. Slipenyuk, V. Kuprin, Y. Milman, V. Goncharuk, J. Eckert, Properties of P/M processed particle reinforced metal matrixcomposites specified by reinforcement concentration and matrix-to-reinforcement particle siz ratio, Acta Mater. 54 (2006) 157-166. https://doi.org/10.1016/j.actamat.2005.08.036 [6] F. Tang, M. Hagiwara, J.M. Schoenung, Microstructure and tensile properties of bulk nanostructured Al-5083/SiCp composites prepared by cryomilling, Mater. Sci. Eng. A 407 (2005) 306-314. https://doi.org/10.1016/j.msea.2005.07.056 [7] K.B. Lee, H. Kwon, Strength of Al-Zn-Mg-Cu matrix composite reinforced with SiC particles,

Metall.

Mater.

Trans.

A

33

(2002)

455-465.

https://doi.org/10.1007/s11661-002-0106-5 [8] A. Kalkanli, S. Yilmaz, Synthesis and characterization of aluminum alloy 7075 reinforced with

silicon

carbide

particulates,

Mater.

Des.

29

(2008)

775-780.

https://doi.org/10.1016/j.matdes.2007.01.007 [9] M.M. Sharma, M.F. Amateau, T.J. Eden, Aging response of Al-Zn-Mg-Cu spray formed 27

alloys and their metal matrix composites, Mater. Sci. Eng. A 424 (2006) 87-96. https://doi.org/10.1016/j.msea.2006.02.047 [10] K.H. Min, B.H. Lee, S.Y. Chang, Y.D. Kim, Mechanical properties of sintered 7xxx series AI/SiCp

composites,

Mater.

Lett.

61

(2007)

2544-2546.

https://doi.org/10.1016/j.matlet.2006.09.062 [11] R. Dasgupta, H. Meenai, SiC particulate dispersed composites of an Al-Zn-Mg-Cu alloy: Property comparison with parent alloy, Mater. Charact. 54 (2005) 438-445. https://doi.org/10.1016/j.matchar.2005.01.012 [12] N.V.R. Kumar, E.S. Dwarakadasa, Effect of matrix strength on the mechanical properties of

Al-Zn-Mg/SiCp

composites,

Compos.

Part

A

31

(2000)

1139-1145.

https://doi.org/10.1016/S1359-835X(00)00062-2 [13] M. Manoharan, J.J. Lewandowski, Effect of reinforcement size and matrix microstructure on the fracture properties of an aluminum metal matrix composite, Mater. Sci. Eng. A 150 (1992) 179-186. https://doi.org/10.1016/0921-5093(92)90110-M [14] S.I. Hong, G.T. Gray, Microstructure and microchemistry of an Al-Zn-Mg-Cu alloy matrix-20 vol.% SiC composite, Acta Metall. Mater. 40 (1992) 3299-3315. https://doi.org/10.1016/0956-7151(92)90043-E [15] Z.L. Shi, S. Ochiai, M. Hojo, J. Lee, M.Y. Gu, H. Lee, R.J. Wu, The oxidation of SiC particles and its interfacial characteristics in Al-matrix composite, J. Mater. Sci. 36 (2001) 2441-2449. https://doi.org/10.1023/A:1017977931250 [16] K.B. Lee, H.S. Sim, S.Y. Cho, H. Kwon, Reaction products of Al-Mg/B4C composite 28

fabricated by pressureless infiltration technique, Mater. Sci. Eng. A 302 (2001) 227-234. https://doi.org/10.1016/S0921-5093(00)01831-1 [17] W.C. Ma, J.L. Gu, Y. Zhang, M.M. Wang, Effect of SiC particles on ageing behaviour of SiCp/7075 composites, J. Mater. Sci. Lett. 16 (1997) 1867-1869. [18] X.J. Wang, X.S. Hu, K. Wu, L.Y. Wang, Y.D. Huang, Evolutions of microstructure and mechanical properties for SiCp/AZ91 composites with different particle contents during extrusion,

Mater.

Sci.

Eng.

A

636

(2015)

138-147.

https://doi.org/10.1016/j.msea.2015.03.058 [19] W.S. Miller, F.J. Humphreys, Strengthening mechanisms in particulate metal matrix composites, Scr. Metall. Mater. 25 (1991) 33-38. [20] Y.Z. Li, Q.Z. Wang, W.G. Wang, B.L. Xiao, Z.Y. Ma, Interfacial reaction mechanism between matrix and reinforcement in B4C/6061Al composites, Mater. Chem. Phys. 154 (2015) 107-117. https://doi.org/10.1016/j.matchemphys.2015.01.052 [21] Q. Zhang, B.L. Xiao, Z.Y. Liu, Z.Y. Ma, Microstructure evolution and elemental diffusion of SiCp/Al-Cu-Mg composites prepared from elemental powder during hot pressing, J. Mater. Sci. 46 (2011) 6783-6793. https://doi.org/10.1007/s10853-011-5636-z [22] T. Marlaud, A. Deschamps, F. Bley, W. Lefebvre, B. Baroux, Influence of alloy composition and heat treatment on precipitate composition in Al-Zn-Mg-Cu alloys, Acta Mater. 58 (2010) 248-260. https://doi.org/10.1016/j.actamat.2009.09.003 [23] A. Urena, E.E. Martinez, P. Rodrigo, L. Gil, Oxidation treatments for SiC particles used as reinforcement in aluminium matrix composites, Compos. Sci. Technol. 64 (2004) 29

1843-1854. https://doi.org/10.1016/j.compscitech.2004.01.010 [24] G.E. Kiourtsidis, S.A. Skolianos, G.A. Litsardakis, Aging response of aluminium alloy 2024/silicon carbide particles (SiCp) composites, Mater. Sci. Eng. A 382 (2004) 351-361. https://doi.org/10.1016/j.msea.2004.05.021 [25] M Strangwood, C.A. Hippsley, J.J. Lewandowski, Segregation to SiC/Al interfaces in Al based

metal

matrix

composites,

Scr.

Metall.

Mater.

24

(1990) 1483-1487.

https://doi.org/10.1016/0956-716X(90)90418-G [26] K. Kondoh, H. Fukuda, J. Umeda, H. Imai, B. Fugetsu, Microstructural and mechanical behavior of multi-walled carbon nanotubes reinforced Al-Mg-Si alloy composites in aging treatment, Carbon 72 (2014)15-21. https://doi.org/10.1016/j.carbon.2014.01.013 [27] B. Cai, B.L. Adams, T.W. Nelson, Relation between precipitate-free zone width and grain boundary

type

in

7075-T7

Al

alloy,

Acta

Mater.

55

(2007)

1543-1553.

https://doi.org/10.1016/j.actamat.2006.10.015 [28] K.T. Aust, R.E. Hanneman, P. Niessen, J.H. Westbrook, Solute induced hardening near grain boundaries in zone refined metals, Acta Metall. 16 (1968) 291-302. https://doi.org/10.1016/0001-6160(68)90014-X [29] V.C. Nardone, K.M. Prewo, On the strength of discontinuous silicon-carbide reinforced aluminum

composites,

Scr.

Metall.

20

(1986)

43-48.

https://doi.org/10.1016/0036-9748(86)90210-3 [30] R.J. Arsenault, N. Shi, Dislocation generation due to differences between the coefficients of

thermal-expansion,

Mater.

Sci. 30

Eng.

81

(1986)

175-187.

https://doi.org/10.1016/0025-5416(86)90261-2 [31] F. Tang, I.E. Anderson, T. Gnaupel-Herold, H. Prask, Pure Al matrix composites produced by vacuum hot pressing: tensile properties and strengthening mechanisms, Mater. Sci. Eng. A 383 (2004) 362-373. https://doi.org/10.1016/j.msea.2004.05.081 [32] K. Ma, H. Wen, T. Hu, T.D. Topping, D. Isheim, D.N. Seidman, E.J. Lavernia, J.M. Schoenung, Mechanical behavior and strengthening mechanisms in ultrafine grain precipitation-strengthened

aluminum

alloy,

Acta

Mater.

62

(2014)

141-155.

https://doi.org/10.1016/j.actamat.2013.09.042 [33] R.L. Fleischer, Solution hardening by tetragonal distortions-application to irradiation hardening

in

fcc

crystals,

Acta

Metall.

10

(1962)

835.

https://doi.org/10.1016/0001-6160(62)90098-6 [34] Z. Zhang, D.L. Chen, Consideration of Orowan strengthening effect in particulate reinforced metal matrix nanocomposites: A model for predicting their yield strength, Scr. Mater. 54 (2006) 1321-1326. https://doi.org/10.1016/j.scriptamat.2005.12.017 [35] H. Wen, T.D. Topping, D. Isheim, D.N. Seidman, E.J. Lavernia, Strengthening mechanisms in a high-strength bulk nanostructured Cu-Zn-Al alloy processed via cryomilling and spark plasma sintering, Acta Mater. 61 (2013) 2769-2782. https://doi.org/10.1016/j.actamat.2012.09.036 [36] O.B. Bembalge, S.K. Panigrahi, Development and strengthening mechanisms of bulk ultrafine grained AA6063-SiC composite sheets with varying reinforcement size ranging from

nano

to

micro

domain,

J.

Alloys 31

Compd.

766

(2018)

355-372.

https://doi.org/10.1016/j.jallcom.2018.06.306 [37] R. Vogt, Z. Zhang, Y. Li, M. Bonds, N.D. Browning, E.J. Lavernia, J.M. Schoenung, The absence of thermal expansion mismatch strengthening in nanostructured metal-matrix composites,

Scr.

Mater.

61

(2009)

1052-1055.

https://doi.org/10.1016/j.scriptamat.2009.08.025 [38] J.C. Shao, B.L. Xiao, Q.Z. Wang, Z.Y. Ma, K. Yang, An enhanced FEM model for particle size dependent flow strengthening and interface damage in particle reinforced metal

matrix

composites,

Compos.

Sci.

Technol.

71

(2011)

39-45.

https://doi.org/10.1016/j.compscitech.2010.09.014 [39] Z.Y. Liu, Q.Z. Wang, B.L. Xiao, Z.Y. Ma, Clustering model on the tensile strength of PM processed

SiCp/Al

composites,

Compos.

Part

A

41

(2010)

1686-1692.

https://doi.org/10.1016/j.compositesa.2010.08.007 [40] X.L. Guo, Q. Guo, J.H. Nie, Z.Y. Liu, Z.Q. Li, G.L. Fan, D.B. Xiong, Y.S. Su, J.Z. Fan, D. Zhang, Particle size effect on the interfacial properties of SiC particle-reinforced Al-Cu-Mg

composites,

Mater.

Sci.

https://doi.org/10.1016/j.msea.2017.11.068

32

Eng.

A

711

(2018)

643-649.

Highlights SiC particles intitiated weakening effect on tensile properties of SiC/Al-Zn-Mg-Cu composites Precipitation behaviors of η′-MgZn2 phases and alloy elements distribution were analyzed High-strength SiC/Al-Zn-Mg-Cu composites were fabricated by adjusting Mg content

Conflict of interest statement We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work. There is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled “An investigation on particle weakening in T6-treated SiC/Al-Zn-Mg-Cu composites”.