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Advanced Composites and Hybrid Materials

, Volume 2, Issue 3, pp 549–570 | Cite as

Hot rolling of MWCNTs reinforced Al matrix composites produced via spark plasma sintering

  • Behzad Sadeghi
  • Pasquale CavaliereEmail author
  • Ghasem Azimi Roeen
  • Martin Nosko
  • Morteza Shamanian
  • Veronika Trembošová
  • Štefan Nagy
  • Niloofar Ebrahimzadeh
Original Research
  • 382 Downloads

Abstract

Aluminum/CNT nanocomposite sheets, with appropriate dispersion and interfacial bonding, were fabricated by a combination of powder metallurgy, spark plasma sintering (SPS), and hot rolling. The effects of CNT content as well as plastic deformation, on the microstructure and mechanical properties of the obtained nanocomposite, were investigated. The composite reinforced by 0.5 wt.% CNTs showed an optimal dispersion of CNTs into the aluminum matrix after both SPS and hot rolling. Minimum CNT damage and minimum carbide formation were observed after hot rolling. The best comprehensive mechanical properties corresponded to the sheet of Al-0.5 wt.% CNT nanocomposite thanks to the strong interfacial bonding between Al and CNTs, full densification of the nanocomposites as well as the uniform dispersion of the CNTs into the aluminum matrix. Hardness measurements showed that the maximum hardness was obtained for sheets containing 1.5 wt.% CNTs in both the as-SPS and the as-hot rolled conditions. Load transfer, Orowan, and grain size strengthening mechanisms could affect the increase of strength as well as the combination of strength and ductility of the sheets of Al-CNT nanocomposites.

Graphical abstract

Aluminum/CNT nanocomposites were hot rolled without reinforcing damage. The optimal dispersion of 1.5% CNTs led to increased mechanical properties.

Keywords

CNT SPS Hot rolling Strengthening mechanisms 

1 Introduction

Carbon nanotubes (CNTs) have attracted the scientific and industrial attention in the recent past because of their incredible elastic modulus and strength. The continuous interest in the very promising properties of nanocomposites led to the development of many types of metal matrix composites reinforced with CNTs [1, 2, 3, 4, 5]. Obviously, the main processing difficulty is represented by the uniform dispersion of the reinforcements into the matrix without damaging the nanotubes. Here, it is necessary to avoid agglomeration in order to optimize the grain boundary pinning effect of the reinforcements [6, 7]. Among the different processing routes available to produce CNT-reinforced aluminum matrix composites, the powder metallurgy technology, in particular Spark plasma sintering (SPS), appears very promising for the production of sounds components with optimal reinforcement distribution by lowering damaging [8, 9]. First of all CNTs allow for the formation of networks around the sintering particles acting as grain boundary pinning [10]. Normally, the SPS is hard to obtain a strong bonding strength between the Al powder particles and completely eliminate the residual pores due to the short diffusion time, so it is necessary to strengthen the interfacial bonding and enhance the densification by post-sintering deformation. Additionally, one of the main limitations is due to the small aluminum particles limiting the reinforcement percentage in this kind of composites [11, 12]. Another important issue is represented by the uniform dispersion of the nanotubes that can be improved by plastic deformation after powder metallurgy processing [13, 14, 15, 16]. To overcome these drawbacks, relentless quests have been done. The post-processing could improve the interfacial bonding through the elimination of the residual pores, concurrent improvement of Al-Al grains and Al-CNT interfacial bonding, breaking of the CNT clusters, and high level of alignment of CNTs [13, 15, 17]. It should be emphasized that the porosity has a significant effect on the mechanical properties. Therefore, it is necessary to strengthen the interfacial bonding and enhance the densification by post deformation. For example, Chen et al. [12, 18] as well as Kown et al. [16, 18] used hot extrusion as post-sintering processing to improve the bonding strength between the Al powder particles and the dispersion uniformity of the CNTs in the Al matrix. Zare et al. [19, 20, 21] applied equal channel angular pressing (ECAP) as post-sintering processing to consolidate the fabricated materials. The well-densified composites with only 2 vol.% CNTs after eight ECAP passes exhibited approximately 30% of increase in the yield strength compared to the pure Al samples. Considering the increasing requirement of the flat products of Al and Al-based composites instead of the bar products provided by hot extrusion or ECAP, it is necessary to pay attention to the microstructures and mechanical properties of the flat products of the Al/CNTs composites. The flat products of Al-CNT nanocomposites not only provide us an excellent opportunity to study the nature of solid interfaces (Al-CNT interfacial bonding as well alignment of CNTs) and the extension of our understanding of the structure–property relationship in nanocomposite materials down to the nanometer or ultrafine grain regime but also present an attractive potential for technological applications with their novel properties. Microstructure control during hot rolling of aluminum sheet was somewhat limited and was controlled by process parameters such as roll speed, temperature and reduction per pass. Normally, plastic deformation such as rolling, extrusion, and drawing will form a strong texture, particularly at room temperature. In contrast, during hot rolling process, the texture and microstructure evolutions as a result of dynamic recovery and recrystallization have a remarkably effect on the tensile strength and ductility.

In the present study, flat products of carbon nanotube (CNT)-reinforced pure aluminum (Al) composites were fabricated by spark plasma sintering and hot rolling. The effects of content of the CNTs on microstructural evolution, strengthening mechanism, and fracture mode of the composites are discussed.

2 Experimental procedure

Raw materials including commercial gas atomized Al powder (ECKA granules Japan Co. purity 98.80%) with average particle size of 30 μm and multi-walled carbon nanotubes (MWCNTs) (US nano Com, purity > 99.81%, average diameter of 30 nm, average length 30 μm) were mixed using a planetary ball mill. Energy dispersive X-ray spectroscopy results of the MWCNTs showed the following composition (%): C 99.46, Al 0.19, Cl 0.02, Co 0.09, S 0.24.

The morphologies of the MWCNTs and Al powder particles are shown in Figs. 1 and 2, respectively. Figure 1a, b shows that the outer surface of the CNTs are clean. Figure 2 shows that the shape of aluminum particles is spherical. The aluminum particles dimensions are in a narrow range. Figure 1a depicts that the CNTs length is about 10–30 μm. Figure 1b illustrates that the outer diameter of the CNTs is in the range 2–30 nm.
Fig. 1

MWCNTs aspect

Fig. 2

Aluminum particles morphology

Spark plasma sintering (SPS) was used to produce the nanocomposite materials. Before mixing with the Al powder, the as-received MWCNTs were ultrasonically treated in ethanol for 15 min to decrease the agglomeration. Then the suspension of the CNTs was mixed with Al powder for 7 h by a planetary ball mill at room temperature under argon atmosphere. To prevent severe cold welding, stearic acid (0.2 wt.%) as the process control agent (PCA) was added to the mixture during ball milling. The rotation speed, the ball to powder weight ratio, and other conditions were the same described in our previous works [22, 23]. The mixed powders were heated at 70 °C for 24 h in a vacuum oven to favor the ethanol evaporation. According to previous reports [22, 24, 25, 26, 27], 0.5 and 1.5 wt.% of CNTs were found to be appropriate contents and exhibited reasonable properties enhancement, and thus in this work 0.5 and 1.5 wt.% Al/CNT composite were chosen. In this regard, two kinds of samples were prepared, the first one by adding 0.5% of carbon nanotubes, the second one by adding 1.5% of MWCNTs. For comparison, pure aluminum powders were selected for the SPS process.

Composites were sintered by SPS technique using Model 10 Ton (Nanozint 10) equipment manufactured by KHPF Inc., Iran [22, 23, 28, 29, 30]. The desired amount of milled powder blend (approximately 15 g) was loaded into a graphite die of 40 mm in diameter. The graphite die was ISG-M303 with a thermocouple hole drilled into the punch. Sintering was conducted under very low pressure (approximately 0.0075 MPa), and the heating and cooling rates were 75 and 30 °C min-1, respectively. Pulsed DC current, with a 36-ms on, 8 ms off pulsed profile was applied. After reaching the chamber internal pressure, the furnace cavity is filled with Argon. The gas removal takes place during the initial stages of heating in SPS [31]. The mixture powders and mold were heated from room temperature to the sintering one of 550 °C, with heating rate and holding time showed in Fig. 3. The whole sintering period (including heating and cooling time) was less than 30 min. Samples were then furnace cooled to room temperature in a vacuum atmosphere. The achieved composite disks were 40 mm in diameter and 5 mm in height. The as-SPSed samples were heated in a furnace at 500 for 2 h and then hot rolled using a laboratory rolling mill with a loading capacity of 20 tons. Hot rolling was done up to 50% reduction in thickness at 400 °C. The roll diameter was 125 mm and the rolling speed was set at 4 m/min.
Fig. 3

Temperature and pressure variations during SPS

The oxygen contamination content of the raw aluminum powder, the Al-CNT mixed powders for all the studied conditions, was measured by LECO (TC-444, LECO Corp., USA).

X-ray diffractometer (XRD) was used to identify the existing phases of samples before and after SPS process. A Philips diffractometer (40 kV) with Cu Kα radiation with λ = 0.15406 nm over 20–100° 2θ was used for measurements. XRD scans were performed with a step size of 0.05° and a dwell time per step of 2 s.

The grain size of Al was estimated from the broadening of XRD peaks using the Williamson-Hall equation, Eq. 1
$$ {\beta}_s\cos \theta =\frac{K\lambda}{d}+2\varepsilon\ \sin \theta $$
(1)
where βs is full-width at half-maximum of the diffraction peak, θ is the diffraction angle, λ is the X-ray wavelength, d is the crystallite size, and ε is the lattice strain.

The microstructural analysis was performed by employing a field emission-scanning electron microscopy (SEM, Jeol JSM-7600F) equipped with an EBSD system (Oxford NordlysNano EBSD, UK) operated at 15 kV and current of approximately 12 nA, in order to assess the microtextures, grain orientations, and grain boundary character distribution (GBCD). Meanwhile, data evaluation at the mid-section of the specimens was collected using the Resmat software. Data processing was then carried out using Mambo and Tango software in order to define low-angle grain boundaries (LAGB, angle range between 2° and 15°), HAGB (angle higher than 15°), orientation distribution function (ODF), and pole figures (PFs) as well misorientation distribution curves. However, it is worth noting that some unindexed regions remained within the investigated area pores, grain boundaries and unidentified phases were replaced using standard noise reduction post process attached in software by way of copies of neighboring points. In our study, MAD values were less than 0.5°, which resulted to 98% indexing for all the scans. Metallographic samples were carried out according to the common procedure of grinding and mechanical polishing. The final polishing step was performed by Oxide polishing Suspension-Type S (OP-S) 0.04 μm. OP-S is a chemically aggressive suspension but is based on colloidal silica. OP-S has a pH of 9.8 and is especially suited for polishing of very ductile materials such as aluminum.

The microstructures of the samples were also characterized using transmission electron microscopy (Jeol TEM 1200EX, FEI STEM Titan Themis). Samples were sectioned by EDM in a direction parallel to the applied pressure. Thin foils for TEM studies were prepared using mechanical grinding followed by ion milling with a GATAN PIPS II machine.

Raman spectroscopic measurements (SOLAR TII Nanofinder, Tokyo Instruments Co. Ltd., Japan) allowed to determine the ID/IG ratio depending on the CNTs after each process.

The composites density was measured following the ASTM B962 - 14. After polishing, the specimens were first weighed in air and then after immersion in distilled water [28, 32]. The density values were obtained by averaging five measurements per sample. To accurately compare the extent of densification that occurred during sintering, a theoretical density was calculated for each powder mixture based on the rule of mixtures approach [22, 28, 29]. A statistical analysis of the particle size and porosity distributions of SPSed specimens was performed by optical microscopy and image analysis.

Mechanical properties were determined by microindentation and tensile testing at room temperature. The specimens for uniaxial tensile tests were prepared by employing elecrodischarge machining (EDM) according to ASTM E8 for small size specimens, with a gauge length of 6.4 mm, 25 mm overall length and 4 mm thickness [22, 30, 31].

The microhardness, using the Rockwell H scale, was carried out using an Akashi apparatus with a 0.3-kg load for 15 s. The tests were performed in load control. The microhardness measurements were done on seven randomly selected points for each specimen to obtain an average value of hardness.

The specimens for uniaxial tensile tests were prepared by employing electro discharge machining (EDM) according to ASTM E8 for small size specimens. The tensile tests were conducted at room temperature with an initial strain rate of 10–4 s−1 using a Hounsfield H50KS machine.

3 Results and discussion

3.1 X-ray diffraction analyses

Figure 4 depicts X-ray diffraction patterns of the SPSed Al-MWCNT nanocomposites with different nanotube contents (0.5 and 1.5 wt.%). For comparison, the XRD pattern of pure aluminum is plotted.
Fig. 4

XRD spectra of the SPSed composites and pure Al

No peaks corresponding to Al4C3 phase (usually seen at about 55°) can be sighted from the XRD patterns for different MWCNT content. The XRD pattern of the pure aluminum reveals only the aluminum diffraction peaks. Many works [33, 34] have used the ultrasonication as an effective technique to de-agglomerate and subsequently to obtain uniform dispersion of CNTs in different matrices. Simoes et al. [33] concluded that the ultrasonic dispersion of the CNTs leads to agglomeration and degradation after 15 min. It is known that Al4C3 particles are nanometric, forming in localized area. Additionally, they result from the interfacial reaction between the damaged CNTs and the aluminum matrix. The supplying source of the carbon, due to Al4C3 formation, is the amorphous carbon or the defective region of CNTs such as their walls. The production processes promote the formation of the carbide. However, these carbides are identified only through high-resolution transmission electron microscopy (HRTEM). Therefore, it is reasonable that no peaks of Al4C3 are detected by X-ray diffraction. However, we expected that in TEM pictures (given in the following) Al4C3 could be observed. This aspect is consistent with other similar studies [18, 33]; it can be concluded that only a barely detectable amount of Al4C3 formed in the composites at the employed sintering temperature. However, it can be concluded that the formation of aluminum carbide crystals during SPS strongly depends on the crystallinity and purity of MWCNTs dispersed within the aluminum matrix [35]. The purity of the used MWCNTs in our study was the same measured in [35].

The possibility of the structural degradation of the CNTs is evaluated through Raman spectroscopy investigations of the hot rolled nanocomposites. Casiraghi et al. [36] showed that first-order peaks of carbon materials lie at 1560 and 1360 cm−1 which are typical for MWCNTs [36, 37]. These peaks were called G and D peaks, respectively. It is known that G peak is due to vibrational modes of atoms in both rings and chains while the D peak is due to the breathing modes of sp2 atoms in the rings [36]. The G and D bond positions are sensitive to graphene sheet and structural defects, respectively. The presence of the D bond significantly reflects that there are a lot of defects in the materials. Generally, the intensity ratio of the D and G bands (ID/IG) is related to the quality of the internal CNTs or to the degree of damage of CNTs [38]. An increase in this ratio depicts an evolution in the structural defects in the CNTS while a decrease indicates a change in the structure of the CNTs.

Figure 5 reveals the Raman spectra of the MWCNTs and the hot rolled nanocomposites. The initial MWCNTs show both D and G bonds as previously reported by Keszler et al. [39]. The milling process leads to the degradation of some of CNT walls and thus resulting in the formation of the defects in the CNTs structure. It is known that the presence of the damages both in tip and structure of CNTs leads to the reaction between the Al and the defective CNTs and subsequently to the formation of Al4C3. On the other hand, the aforementioned defects caused by milling as well the infiltration of Al atoms in CNTs chains lead to distortion of sp2 bonding structure and subsequently the G peak shift toward higher wave number. On the other hand, the imposed strain during ball milling operation leads to an increase in the inter-atomic distance of carbon-carbon in the CNTs. It is worth to note that no shift in the G band occurs during the subsequent hot rolling process. However, no clear D peak shift was observed. The findings were consistent to previously reported results [12, 34, 38, 40, 41].
Fig. 5

Raman spectra of the hot rolled nanocomposites and as-received MWCNTs

Besides, the ID/IG ratio that represents the degree of structural defects in the CNTs, regardless of the type of production process, increased by increasing the CNTs concentration as given in Table 1. The peak intensity ratio of the initial CNTs WAS 0.67, increased to 0.92 and 1.12 after ball milling for 0.5 and 1.5 wt.% CNTs. However, this ratio showed limited variation (1.08 and 1.23) after hot rolling. Therefore, it can be concluded that the ball milling is a severe plastic deformation process which introduces some damages such as bending and breaking of the CNTs, amorphous impurities, and even open tips of CNTs. Further CNT damages did not occur during the rolling process because of its high elastic modulus. It is well-know that the formation of Al4C3 occurred through two routes. The first route is based on the reaction of Al with carbon at T ≥ 500°C, and the other one is based on the diffusion of carbon source into the aluminum oxide skins at temperatures higher than 2000 °C [42, 43]. The using of the graphite die damaged CNTs and even the using of the process control agent (PCA) [29]. Due to the existence of the disordered parts, defects and open tips of the CNTs as well sintering temperature higher than 500 °C Al4C3 formed. Furthermore, the generation of local high temperatures onto aluminum surface particles during SPS process assists the diffusion of carbon, resulting in the formation of the Al4C3.
Table 1

Influence of the CNT concentration and process temperature on ID/IG ratio

Material

ID/IG

Initial MWCNT

0.67

Al-0.5 CNT milled powders

1.09

Al-0.5 CNT as-hot rolled

1.11

Al-1.5 CNT milled powders

1.15

Al-1.5 CNT as-hot rolling

1.21

It can be attributed to the residual stress release resulted by milling process that occurs in the SPS process. The hot rolling process conducted at high temperature for the understudy nanocomposites (400 °C), stimulates the release of residual stress through provoking accumulated dislocations back of the obstacles such as CNT, probable nanocarbide particles of the Al4C3 as well as grain boundaries (GBs). In this circumstance, reducing the system energy through annihilation of the dislocations climb occurred. The variations of the ID/IG ratio justified this mentioned mechanism.

It is reported that the formation of Al4C3 phase in Al-CNT nanocomposites, that mainly locates in the CNTs walls, can be the cause of the slight increase in the ID/IG ratio [33, 34]. Therefore, G peak shift amount as well as the ID/IG ratio changed. Meanwhile, it can be clearly seen that the ID/IG ratio increased (~1.82 % and 5.3 % increase for 0.5 and 1.5 wt. % of CNT) after hot rolling process. It can be attributed to the high formation of the aluminum carbide in higher concentrations of damaged CNTs, bonding condition and even the infiltration of Al atoms in CNTs. It is important to note the increasing of the ID/IG ratio cannot be attributed to the structural changes of the CNTs. In fact, the most important reason for structural change in a CNT is the reaction of the CNT with oxygen at high temperature. However, the presence of the Al alongside CNTs, (aluminum oxide have a higher priority than CNT oxide up to about 2000 °C, with reference to the Ellingham diagram [44]) impedes the decomposition of the CNTs. So, the trend in the variation of the ID/IG ratio is probably due to the release of the residence stress of the CNTs caused by hot rolling. Therefore, it can be concluded that MWCNTs still retain considerably their high crystalline integrity while the initial milling, SPS and then hot rolling process have caused some damages to the pristine MWCNT surface and tips [45, 46].

Figure 6 depicts the microstructure of Al-CNT nanocomposites before and after hot rolling. Figure 6a–d shows that the hot rolled nanocomposites have the better distribution of CNTs comparing to the as-SPSed material. CNTs are preferentially located as clusters in the grain boundaries (Fig. 6b inset). While it is guessed that the hot rolling process leads to the breaking of the CNT clusters and thus CNTs are embedded into the aluminum matrix (hereafter shown by the black arrows). The shear and compression stresses caused by hot rolling lead to breaking the CNTs agglomerates (hereafter shown by red elliptic), subsequently to their uniform distribution in the aluminum matrix. Additionally, the aluminum carbide particles (hereafter shown by yellow arrow) probably are broken and are distributed in the aluminum matrix as well. Moreover, it seems that the mean grain size of the aluminum for as-hot rolled samples is finer than that of the as-SPSed samples. Figure 6e, f shows that the aluminum grains, as well as the CNTs, are elongated and aligned along the rolling direction. However, the Al4C3 increased for the as-hot rolled Al-1.5wt.%. As can be seen in Fig. 6a–d, by increasing the CNT content, the agglomeration for the as-SPSed samples and the presence of Al4C3 for the as-hot rolled samples increase. For higher CNT content, the damage of the CNTs is more pronounced, resulting in larger reaction with the aluminum particles. In such cases, the formation of aluminum carbide is favored.
Fig. 6

SEM of a, b as-SPSed, inset higher magnification, c, d as-hot rolled, e, f high magnification of as-rolled nanocomposites. a, c, e Al-0.5 wt.% CNT and b, d, f Al-1.5 wt.% CNT

TEM images of the SPSed Al-1.5 wt.% CNT nanocomposite were shown in Fig. 7. By a close inspection to these figures, it is clearly revealed the area of Al4C3 crystals (confirmed with SAED pattern, Fig. 7b inset) in as-SPSed Al-1.5 wt.% CNT nanocomposite. Hexagonal aluminum carbide Al4C3 presence was confirmed with ring SAED pattern. The calculated inter-planar distances matched with Al4C3 PDF card number 00-050-0740. The formation of carbide will be beneficial for a strong bonding between the aluminum matrix and CNTs [18]. However, it affected the wetting angle of Al and CNT. It is reported that the formation of interfacial aluminum carbide resulted in the reduction of the wetting angle of Al/CNT by 55° [47]. Although, the aluminum carbide is obtained as a result of the reaction between damaged CNTs and aluminum [48].
Fig. 7

TEM micrographs a low and b high; magnification of as-SPSed Al-1.5 wt.% CNT nanocomposite

It can be seen that the formed Al4C3 microcrystals have both dumbbell and are tube type which are attributed to the chemical reaction between aluminum with the tip of a CNT and defective CNTs, respectively [15]. On the other hand, the aluminum carbide formed at the aluminum/CNT interface mainly consisted of some of the CNTs walls or the entire CNTs. These carbides are usually nanometric and form in localized areas. The formation of Al4C3 plays a significant role in the stress transfer from the aluminum matrix to the CNTs, strongly affecting the mechanical properties of the Al-CNT nanocomposites. Major CNTs embedded in the aluminum matrix result in a uniform dispersion promoting work hardening rate and improved mechanical properties. Some aluminum oxide skin inside the aluminum matrix, despite the surface cleaning effect of the SPS process, is shown by red arrows in all figures. These phases were identified by EDS, the SAD pattern, and STEM.

SEM-EDX map representing the oxide skin on the aluminum grains is shown in Fig. 8. Amorphous carbon, aluminum oxide, graphite, and aluminum carbide in both samples were identified. Aluminum oxide consists of both amorphous and crystalline forms [49]. The presence of the broken Al2O3 particles, CNTs as well as aluminum carbide particles into aluminum matrix generally leads to improvement of tensile strength. It has been specified that the thin oxide film was easily formed on the surface of aluminum powder particles which affects the aluminum atom inter-diffusion between the adjacent aluminum particles.
Fig. 8

a SPSed sample. b Oxygen EDX map representing the aluminum oxide layer on the aluminum powders

Figure 9 shows the bright field TEM image of Al-CNT nanocomposite after hot rolling. Figure 9a–d shows that CNT agglomerates are broken, consequently attaining a uniform distribution into the aluminum matrix. However, it can be seen a few amount of aluminum oxide and carbide aluminum into aluminum matrix as well. In fact, by increasing the CNT content, both size and quantity of the individual and clusters of the CNTs increased. The elongation and tensile strength of the Al-CNT nanocomposites are mainly affected by the following factors: matrix ductility and strength, dispersion of CNTs, interface reaction, and damage of CNTs. Figure 9b, d shows that CNTs are embedded into aluminum particles matrix. Moreover, it seems that the individual CNTs are aligned along with rolling direction. The CNTs were straightened and aligned along the rolling direction after hot rolling (Fig. 9b, d). Among effectible factors, the CNT dispersion into aluminum matrix has a significant influence on both strength and ductility.
Fig. 9

TEM micrographs of as-hot rolled a, b Al-0.5 wt.% CNT and c, d Al-1.5 wt.% CNT nanocomposite

To enhance the ductility of Al-CNT nanocomposites without dramatic loss of their strength, it is important to hinder both plastic strain instability and crack nucleation instabilities. The presence of stored strain energy caused by both the primary ball milling and hot rolling process, besides high temperature of both SPS and rolling process provided an excellent condition for occurring the recovery and recrystallization for aluminum grains. However, the presence of Al4C3 and Al2O3 particles has an effective role in limiting the aluminum grain growth. By increasing the CNT content, the crystallite size of aluminum decreased. On one hand, the dislocation slip is dominant in aluminum grains, and on the other hand, the settlement of the CNTs into aluminum grains as effective obstacles to any dislocation movement leads to increase the strain hardening. Therefore, the uniform distribution of CNTs into the aluminum matrix in Al-CNT nanocomposites resulted in locking the dislocations movements and thereby intensively accumulated at Al-CNT interfaces. The stress fields of such locked and tangled dislocations contribute to strain hardening. By increasing the locked dislocations quantity for both Al-CNT nanocomposites, after applying the hot rolling process, the strain hardening increased. Asgharizadeh et al., for aluminum matrix nanocomposite containing 3 vol.% of CNTs, fabricated via ball milling and HPT consolidation route and reported that the low ductility is recorded as a result of CNT agglomerations which act as stress concentrators initiating intensive strain localization and crack nucleation [50]. In the current research, the TEM figures showed that for the lower content of CNTs (0.5 wt.%), the most of CNTs are embedded into aluminum grains and apparently are aligned with the rolling directional. However, the formation of some little quantity of aluminum carbide is useful for obtaining a tight and clean Al-CNT interface and consequently the bonding strength between the reinforcement and the matrix as shown in Fig. 10.
Fig. 10

HRTEM images of the Al-0.5 wt.% CNT nanocomposite subjected to hot rolling process showing a the CNT, aluminum carbide and Al matrix, and b the Al/CNT interface

Previously, it has been clarified through Raman spectroscopy analysis that during the hot rolling process probably some damages to the pristine MWCNT surface and tips are revealed. In this regard, in order to specify these defects, HRTEM was employed. Nanocrystalline aluminum carbide phase, adjacent to the CNTs-aluminum particles, were observed. Additionally, the interlayer spacing values corresponding to aluminum carbide as well CNTs were calculated. The value of 0.34 nm is identified as the interlayer spacing of embedded CNTs into the aluminum grain. While, the aluminum carbide particles are identified by the 0.282 nm inter-planar spacing of (012) planes. In order to more and closer understanding of the relationship between the crystallographic planes and inter-planar spacing of Al, CNT, and Al4C3, Table 2 is given.
Table 2

(hkl) versus inter-planar spacing of Al, CNT, and Al4C3

(hkl)Al

d (nm)

(hkl)Al4C3

d (nm)

(hkl)CNT

d (nm)

111

0.234

003

0.834

(002)

0.34

200

0.202

006

0.417

(100)

0.29

220

0.143

101

0.287

 

311

0.122

012

0.282

222

0.117

009

0.278

The calculated values of crystallite or grain size and strain are given in Table 3. At the first glance, as increasing the CNT concentration, the crystallite size decreases. The strain accordingly increases. The grain size of the achieved nanocomposites was below 100 nm. Therefore, in this study, we can conclude that the produced nanocomposites can be considered as UFG materials with unique mechanical properties that originate from their microstructure.
Table 3

Change in crystallite size and strain with CNT content in the nanocomposites calculated from XRD peak analysis

Wt.% CNT

SPSed

Hot rolled

Crystallite size (nm)

Strain

Crystallite size (nm)

Strain

0

96

1.04E−05

120

0.10E−05

0.5

85

2.12E−05

65

0.48E−05

1.5

80

3.28E−05

42

0.75E−05

Decreasing of internal strain is a consequence of recovery occurred during the hot rolling. It is well-known that the structure of GBs in nanostructured materials, such as UFG ones produced through severe plastic deformation, is strained with high densities of disorderedly arranged GB dislocations [51]. It is worth to note that the density of the dislocations decreases after hot plastic deformation comparing to the cold plastic deformation. The required shear stress for initiating slip and generate significant plastic strain is called critical resolved shear stress (τc). Since this stress has a low sensitivity to temperature [52, 53, 54], and also, the size of the crystallites is below 100nm, it is expected that vacancy climbing, dislocation slip as well as cross-slip occurred during the hot rolling process. Consequently, dynamic recovery and dynamic recrystallization occurred. In addition, it is logical to be stated that the significance of the aforementioned mechanism changes with the variations of the grain size. It is known that the introduction of the defects into nanocomposite structure causes the increasing in the τc, resulting in the difficulties for dislocation sliding.

Moreover, the dislocation density in the presence of the CNTs increases if compared to pure aluminum as shown in Fig. 11. It causes a negligible increase in the Schmid factor related to the sliding systems and consequently increases the required τc and work hardening. This causes significant effects on the microstructure and mechanical properties.
Fig. 11

a Computed dislocation density in the nanocomposite as a function of CNT content. b Geometrically necessary dislocation distribution of the nanocomposites

In our previous work, we have found that during SPS partial dynamic recrystallization occurs [55]. Additionally, it is known that the SPSed nanocomposites contain a bimodal grain structure [56, 57]. In addition, it seems that after hot rolling, the grain boundaries fraction (both low angle and high angle) in the nanocomposite is higher than that of the SPSed nanocomposite. This microstructural change is a result of the dynamic recrystallization.

Since SPS process is a sintering technique with high heating rate and intense joule heating, the grain boundary and volume diffusions are active. Additionally, during hot rolling, strain rates are in the range 100 − 102s−1 at temperatures higher than 0.5 Tm [58].

It is known that diffusion is governed by temperature and can act as an outstanding mechanism in hot rolling. Dynamic recovery as a consequence of climb and cross-slip of dislocations occurs readily during hot working of metals of high stacking fault energy such aluminum [58, 59, 60, 61]. During dynamic recovery, dislocations start to rearrange and form low-angle boundaries as sub-grains develop. On the other hand, the heating rate in the hot rolling is slower than that of SPS process. Since the rolling itself causes the formation of dislocations into the grains, the high-temperature influence leads to the activation of plastic deformation mechanisms such as GB migration/sliding or grain rotation. Hence, it is concluded that the occurrence of the dynamic recrystallization during the hot rolling process leads to the formation of the recrystallized fine grains. Meanwhile, it should be noted that the presence of the CNTs and also their propensity to locate at GBs can hinder the mechanisms which lead to the grains coarsening. In addition, during hot rolling, dynamic recrystallization occurs and subsequently the recrystallized grains easily form. The uniform dispersion of the CNTs in the aluminum matrix, especially along with grain boundaries, leads to impeding the aluminum grain boundaries movement, hampering the grain growth and finally the generation of grains smaller compared to the SPSed ones. It is known that homogenous dispersion has a remarkably effect on the recrystallized grain size [23, 29, 31]. All of these effects lead to increase in the strength and hardness of the as-hot rolled Al-CNT nanocomposites.

By increasing the CNT concentration, the thermal stresses as well as dislocation density increase. By contrast, the grain size of the as-SPSed nanocomposite decreased for 0.5 wt.% CNT and increased for 1.5 wt.%, in comparison to pure aluminum. However, after hot rolling, the grain size of both the nanocomposites decreases as a consequence of the hindering of grain boundaries growth due to the presence of different obstacles. It is attributed to the CNT clusters that situated along with the GBs. The homogenous distribution of the CNTs in the aluminum matrix leads to hinder the grain growth. In addition, the uniform distribution of the CNTs in the aluminum matrix splits the applied stresses to lower stresses, leading to decrease the stress accumulation in back of the obstacles such as CNTs. However, in presence of the CNT clusters, the spoliation of the applied stress decreases and consequently the accumulated stress increases. Moreover, when clustering occurs, the CNTs bonding to each other are governed by van der Waals forces. Consequently, the clusters break and therefore have no valuable role in the increasing of the mechanical properties.

The dislocation densities given in Fig. 11a, were calculated by Eqs. 24 [62]:
$$ {\rho}_D=\frac{3}{{\left({D}_{\mathrm{vol}}\right)}^2} $$
(2)
$$ {\rho}_{\varepsilon }=K\frac{\varepsilon^2}{{\overline{b}}^2} $$
(3)
$$ \rho ={\rho_{\varepsilon }{\rho}_D}^{\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$2$}\right.} $$
(4)
where ρ is the dislocation density, ρε and ρD contribution of strain and crystallite size to dislocation density, K = 6π is a constant, \( \overline{b} \) = \( \frac{a}{\sqrt{2}} \) is the Berger vector of FCC, and a is the lattice parameter of the materials.

The main factors affecting the dislocation density in the study were recognized in the CNT reinforcements, Al4C3 carbide, and the hot rolling effects. In the conditions under study, the increasing of the dislocation densities was attributed to the CNTs and Al4C3 effects on pinning of the GBs. Moreover, the geometrically necessary dislocations (GND) density increases by increasing the CNT content. The formation of the GNDs is as a result of uneven deformation and dislocation slips in order to a kind of dislocations coordination. The GND density is affected by the carbide inclusions, CNT reinforcements as well as severe plastic deformation. Therefore, it can be concluded that the GND density increased through hot rolling.

Casati reported that the GND dislocations originate from different coefficient of thermal expansion and also elastic modulus [63]. In Fig. 10b, the calculated values corresponding to the GND dislocation density showed an increasing trend by increasing the CNT content. It could be due to the CNT influence on the GB movements in the Al matrix. The presence of the CNTs as well as nanosized Al carbide has a significant effect on the dislocations pinning. With regard to the CNTs tending to locate in grain boundaries [18, 64], for the high CNT concentration, the dislocation density increased due to the CNT clusters acting as precipitates during the hot rolling. Owing to the difference in strength between the CNT clusters and the Al matrix, a stress field is generated around the clusters. It causes hindering in the flow of dislocations through the matrix during the hot rolling. However, the high stress field around the CNTs can be relaxed through the generation of GND dislocations, but the high temperature of the hot rolling process facilitates the dislocation climbing and rearranging. In this circumstance, also recovery occurs.

3.2 Mechanical properties

In Al-0.5CNT nanocomposite, due to uniform dispersion of CNTs over the matrix, the created stress field is smaller than that of the CNT clusters. In fact, the stress fields are homogenously distributed throughout the matrix. The induced stress, due to the dislocation pileups in back of the stress concentration points as well as to the applied stress during rolling, leads to the oxide layer of aluminum particles braking and therefore to the intimate Al-Al contact points. It causes increases in the tensile strength as well as in the elastic modulus. The dislocation density, for sample without the CNTs, is clearly reflected in lower amount of strain hardening which is caused by the dislocation pile ups. Therefore, its elongation is higher than that in the samples containing of CNTs. It is known that for high CNT content, the influence of the reinforcements is poor due to agglomeration or clustering. As above mentioned, the clustering can possess a significant role in the generation of the obstacles for the dislocations movement. In addition, the CNTs agglomerates as well as the Al4C3 carbide can also act as Frank–Read source favoring dislocation loops. Increasing the dislocation density, as shown in Fig. 10a, leads to strain hardening. It can be concluded that CNTs play a dual role in strengthening mechanisms depending on their degree of dispersion. Homogenously, dispersion of CNTs increases the elastic modulus for low content of CNTs. While for high content of CNTs, clustering leads to increase in the mechanical properties such as tensile strength and decrease in the elongation and elastic modulus.

The studied materials’ density and porosity after SPS are shown in Fig. 12.
Fig. 12

Density and porosity of the studied materials after SPS

The addition of 0.5% of MWCNTs to pure aluminum leads to an increase in the material density, then it slightly decreases for a percentage of nanotubes of 1.5. This behavior is confirmed by the porosity measurement. Porosity strongly decreases thanks to the carbon nanotubes addition but it reached the lowest value for 0.5%. Even though the oxide layer is covered the surface of the aluminum particles, the relative densities of the SPSed samples were around 98%.

Hot rolling has a general effect on all the materials properties leading to the increase of density if compared to the SPSed condition (Fig. 13).
Fig. 13

Density and porosity of the studied materials after hot rolling

After hot rolling the materials’ density results are very close to the theoretical one for all the nanotubes percentages. Porosity, in this case, increases as increasing the percentage of MWCNTs if compared with the pure aluminum. In addition, it is also worth to note that the CNT content has effect on the density of the as-SPS and as-rolled samples. By comparing the figures given in Figs. 13 and 10, it can be found that by increasing the CNT content, the density is slightly decreased [23, 65]. The Al-0.5 wt.% CNT comparing to the Al-1.5 wt.% CNT has less Al-CNT interfaces and aluminum atom diffusion barriers, so aluminum atoms can readily diffuse, resulting in filling of interstices between the adjacent aluminum particles during both SPS and hot rolling process. Therefore, the presence of high temperature both in SPS and hot rolling process leads to aluminum materials flow with the aid of diffusive mechanisms [23, 66]. It could be influenced by the CNT agglomeration which hinders the plastic deformation of the aluminum matrix around them. However, through hot rolling, some of these agglomerates are broken even in Al-1.5 wt.% CNT nanocomposite and consequently a uniform distribution of CNTs appears into the aluminum matrix. It should be emphasized that some CNT clusters and aluminum carbides formed for Al-1.5 wt.% CNT nanocomposite even after the hot rolling process. One of the key factors in the variation of the density before and after hot rolling is the breaking of alumina layers on aluminum particles as well as CNT clusters during rolling and filling out the pores between particles. This is the main reason for increasing the densities after hot rolling. However, the sintering parameters such as the heating rate, pulsed current, and applied pressure are also responsible for increasing the densities of the produced nanocomposites by SPS [22, 29, 31]. Additionally, many studies under HRTEM analysis depict that the subsequent hot working such as hot rolling or hot extrusion affects the orientation of the CNT within the aluminum matrix [16, 18, 35]. The subsequent hot rolling leads to increase the possibility of the CNTs alignment, consequently the higher densification capacity. Such full densified nanocomposites are the result of the applying of high-pulsed current density (about 2000 A), generating localized microplasma into gaps or at contact points between aluminum particles. It is proved that the microplasma causes the cleaning effect of the powder surface through the generation of the local high temperature and subsequently the evaporation and melting on the surface of the powder particles [67, 68, 69, 70, 71]. Consequently, the diffusivity between aluminum particles even in the presence of the CNTs as reinforcements increases. We also believe that the beneath reasons have a significant contribution to the breakdown or removal of the oxide layer in part.

In fact, the concurrent application of both temperature and pressure (the pressure applied on the punches is usually between 10 and 100 MPa), leads to the reduced cross-section at the particle to particle contacts, the local pressure here is increased by a factor of 100, which results in theoretical local pressures of 1000 to 10,000 MPa. These high pressures can easily break up the oxide layers and create direct metal to metal contacts [13, 18, 47, 63, 64].

The change in the oxygen contamination content for all the samples is listed in Table 4.
Table 4

Oxygen content of the studied samples

Material

Oxygen content (wt.%±2)

Al powder

0.423

As-SPSed pure Al

0.195

As-hot rolled pure Al

0.194

As-SPSed Al-0.5 CNT

0.795

As-hot rolled Al-0.5 CNT

0.791

As-SPSed Al-1.5 CNT

0.778

As-hot rolled Al-1.5 CNT

0.771

The major reason for the decreasing of the oxygen content after the SPS process is the cleaning effect of SPS process. In fact, the pressure caused by the microplasma or spark discharge can effectively eliminate adsorptive gas and any impurities that are present on the surface of the aluminum powder and can likely break up the oxide films on the particle surface [67, 69, 72, 73]. The local pressure concentrates on the particle surfaces and enhances as increasing the sintering temperature [74]. Therefore, at high temperature, the breaking up intensity of the oxide layers increases and subsequently the metal-metal contact points increase. It can be concluded that SPS with the aid of its oxide-cleaning effects lead to decrease in the oxygen content.

Additionally, the weight percent increase of the aluminum leads to higher oxygen content in the nanocomposites irrespective of the CNT content. Furthermore, it is shown no significant change in the oxygen content even after the hot rolling [18]. It is necessary to mention that the reduced atmosphere caused by the using of the carbon mold and graphite sheets during SPS process as well as low pressure into the SPS chamber (70 mbar) contributed to the oxygen content. It is proven that the oxygen content decreased as the SPS temperature increased [74]. Hence, high sintering temperature (about 550 °C) leads to the removal of the aluminum oxide layers on the particles [18, 74, 75]. However, many researchers have identified the presence of the nanosized or ultrafine alumina nanoparticles in the Al-CNT nanocomposites [13, 15, 18]. However, it is reported that the percentage of contribution from alumina particles strengthening is 35.4% in oxygen content of 0.76 wt.% for the Al-1 wt.% CNT nanocomposite [75]. Being the oxygen content in the present study about 0.77–0.79 wt.%, therefore, it is logical to consider the alumina particles strengthening contribution.

The tensile curves of the studied materials after SPS and hot rolling are shown in Fig. 14a. As a general behavior, hot rolling leads to an increase in the tensile strength of the materials. By observing the composites behavior, it is clear how the tensile strength of the 1.5% MWCNT composite is higher with respect to the 0.5% one. The as-SPSed pure aluminum depicted plastic yielding at 45 MPa, while the as-rolled pure aluminum showed plastic yielding at 230 MPa. Additionally, it was shown that all the composites depict a much higher strain hardening than the pure aluminum.
Fig. 14

a Tensile curves of the studied materials. b Relationship between the UTS/YS ratio vs. UTS for the Al-CNTs composites produced by ball milling and subsequent hot rolling

Figure 14b shows that the 0.5 and 1.5 wt.% CNT-aluminum nanocomposites achieve higher UTS/YS ratios (1.14 and 1.15 vs. 1.08) than the pure aluminum. It can be attributed to the good dispersion of the CNTs into aluminum matrix by ultrasonic and ball milling.

As comparing the stress-strain curves, it is found that the elongation to failure decreases by applying hot rolling process. In the same way, this value for pure aluminum decreased after hot rolling. By applying the severe plastic deformation over pure aluminum, the structure consists of ultrafine grains which causes significantly decrease in the strain-hardening capacities. The flow stress had a strong dependence on the temperature and strain rate [76]. During the hot rolling, due to high temperature (higher than 0.5 Tm) the phenomena e.g. dislocation climbing and cross-slip which are thermal activated processes, result in dynamic recovery . Sometimes even partial recrystallization (recrystallization fraction depend on the rolling temperature) occurs [77]. The generated dislocations caused by SPD have organized themselves into cell structures. The evolution of these cells leads to form the low- and high-angle grain boundaries and consequently decrease the dislocation density inside many grains. Therefore, the low ability of strain hardening for pure aluminum is attributed to enhanced dynamic recovery in ultrafine-grained microstructure with large density of grain boundaries relative to the grain volume.

For a constant content of CNTs, the elongation to failure increased up to ~69% and ~50% after hot rolling for 0.5 wt.% and 1.5 wt.% CNTs, respectively. Surprisingly, the elongation to failure of the as-rolled pure aluminum decreased at ~37% as compared to the as-SPSed counterpart. Additionally, the decrement of elongation to failure of the as-rolled nanocomposites with 0.5 wt.% and 1.5 wt.% CNTs, was ~ 32% and ~48% comparing to its pure aluminum counterpart, respectively. Therefore, it can be concluded that the hot rolling and CNT content has a significant effect on the tensile properties.

It is clarified that the introduction of CNTs is one solution for improving the elongation due to block and accumulate the lattice dislocations. In fact, the CNTs located at grain boundaries as well as in the grains interior lead to hindering of the dislocation slip.

In fact, the stress concentration sites such as the alumina particles, the residual CNT clusters, the aluminum carbide as well as the pores, can act as the places for crack nucleation, resulting in the decrease of the elongation. Additionally, during hot rolling which is performed at high temperature and roller pressure, atom to atom bonding fraction between aluminum particles improves. Due to the reduced cross-section at the particle to particle contacts during SPS as well as rolling process, the local pressure increased resulting in theoretical local pressure. This high pressure could simply break up the oxide layers and create metal to metal contacts [31]. In addition, applied pressure during rolling process stimulates and accelerates the breaking up of the CNT clusters, and probably aluminum carbide clusters and therefore increases tensile properties.

On the other hand, the subsequent hot rolling process affects the distribution of the CNTs, aluminum nanocarbides as well as alumina particles within aluminum matrix. Additionally, the applied pressure during hot rolling besides the using the high temperature (more than 0.5 melting temperature) leads to increase the bonding fraction of the aluminum particles. It causes to increase the necking quantity between aluminum to aluminum particles and consequently mechanical properties of the samples increase. It should be also recognized that the formation of aluminum carbide has remarkably effect on the microstructure and mechanical properties of the hot rolled nanocomposites. Kwon et al. [15] reported that the formed carbide were of two different types: one is dumbbell-shaped and the other tube-shaped. It is known that the average size of the Al4C3 was within nanosized (about 20–40 nm [18], about 5 nm in [47]). These aluminum carbides can situate as implant into the aluminum matrix in the hot rolled process. Such implants play an important effect in strengthening through load transfer from the aluminum matrix to the CNTs in the Al-CNT nanocomposites. One the other hand, the formation of the aluminum carbide causes the bonding between aluminum and CNTs to be stronger [47]. In addition, the transfer of any stress in the matrix to the CNTs through the aluminum carbide has been illustrated [15, 18, 78, 79].

The typical stress-strain responses as well the hardness values of the studied materials are summarized in Fig. 15.
Fig. 15

Tensile properties and microhardness of the Al-CNT nanocomposites as function of CNTs percentage

By increasing the CNT content, the as-hot rolled nanocomposites exhibited a tensile strength of around 318 MPa and 352 MPa, which are 74% and 92% higher than that of the pure aluminum, respectively. Moreover, the increment of the tensile strength resulting from hot rolling process was about 44% and 50% for nanocomposites reinforced with 0.5 wt.% and 1.5wt.% of CNTs. The increment in the mechanical properties of the CNT-reinforced aluminum matrix nanocomposites originated from different involved mechanisms such as the load transfer from the matrix to the CNTs [75, 80], Orowan mechanism due to the looping of lattice dislocations at CNTs [22, 50, 75], grain refinement mechanism [50, 81], solid solution strengthening of the carbon atoms [52, 82], particle strengthening induced by the in situ formed or precipitated carbides such as Al4C3 [6, 18, 47], and work hardening mechanism due to dislocations emitted as a consequence of thermal and modulus mismatch [14, 28, 83]. The coupled effects can be theoretically explained via a unified model [84, 85] or the summation of each one contribution. Tensile strength increases by increasing of the MWCNT content to pure aluminum. Figure 15a, showed a significant increase in tensile strength of the nanocomposites, from 266 to 290 MPa for the as-SPSed ones and from 395 to 412 MPa for the as-hot rolled ones. This behavior is confirmed by modulus as shown in Fig. 15c. In addition, a remarkable increase in tensile strength of the as-hot rolled nanocomposites, about 94–102% for 0.5 wt.% and 1.5 wt.% of MWCNTs, respectively, comparing to its pure aluminum counterpart, should be underlined [25]. Here, high CNT percentages lead to an increase in the tensile strength only if high sintering temperatures are employed. In fact, the material strength increases for sintering temperatures higher than 800 850 K. For lower sintering temperatures, a high percentage of nanotubes presented damaging leading to tensile strength drop. Figure 15b shows the hardness of the as-SPSed and as-hot rolled nanocomposites as well one is without the CNTs. The as-hot rolled samples depict much higher hardness comparing to its as-SPSed counterpart. Additionally, the hardness of all samples increases from 33 to 86 HV for as-SPSed and 56–110 HV for as-hot rolled with increasing the MWCNT concentration from 0 to 1.5 wt.%.

In the absence of the CNTs, the strength of the as-hot rolled pure aluminum surprisingly increased (about 120%) comparing to the as-SPSed pure aluminum.

The elongation to failure is dramatically affected by the grain size of nanostructured materials [86, 87]. Ductility is governed by the ability to suppress strain instability associated with macroscopic necking. Plastic strain in most of materials is mediated by the dislocation slip [76]. Therefore, it is expected that tensile ductility in nanocrystalline materials is lower than that of the coarse-grained counterparts. In a review paper [88], the following reasons were identified: the plastic instability, crack nucleation, shear instability, and processing artifacts, e.g., pores [50, 89]. The elongation of the as-hot rolled pure aluminum was about 23% which is about 26% lower than that of its SPSed counterpart. It can be most likely attributed to dynamic recovery and recrystallization that influence both the microstructural changes and the mechanical properties. However, the presence of some undesirable hard inclusions such as aluminum oxide as well as aluminum carbide which could locate at the recrystallized grain boundaries significantly affects the elongation. A minor addition of MWCNTs to the matrix leads to insignificant decrease of the elongation to failure with respect to the pure Al both for the as-SPSed and as-hot rolled samples. In addition, by increasing the CNT content, elongation to failure of the as-hot rolled nanocomposites decreased comparing to its pure aluminum counterpart. Hot rolling leads to the increase of ductility in all samples after SPS process. As it is depicted in Fig. 15d, the decrement of the ductility both in the SPSed and in the hot rolled nanocomposites is significant for a CNT content of 1.5 wt.%. However, the decrement is lower for minor content (0.5 wt.% of CNTs), this was attributed to the uniform distribution of inclusions and CNTs within aluminum matrix. Therefore, it can thus be concluded that both the CNT content and the hot rolling process can significantly affect both tensile properties as well as the hardness for all the samples.

The small grain size of the matrix results in decrease in the elongation to failure [64, 76]. CNTs are often located at GBs where CNTs as well as the carbon nanoinclusions such as Al4C3 result in impeding GB migration and then the grain growth during fabrication process. In contrast, in the sample without CNTs, a substantial grain growth occurs during hot rolling. It is accompanied by a decrease in its strength.

The slop of the true stress-strain curve depicts the strain hardening [90]. All the materials depict a plateau (In the best case a low slop for as-hot rolled nanocomposite containing the 1.5 wt.% of CNT) once starting of plastic deformation. Such behavior reflects that accumulation of the dislocations is lower than its rearrangement and annihilation. Consequently, the dynamic recovery occurs fast. From one side, the simultaneously presence of high temperature and strain during hot rolling process, and on the other, high stacking fault energy of aluminum lead to the formation of both low-angle and high-angle grain boundaries through the dynamic recovery and recrystallization.

In other words, it is supposed that strain hardening has a little effect as strengthening mechanism in the produced nanocomposites. In such conditions, the role of hard reinforcement such as CNTs, Al4C3, and broken Al2O3 layers are significant in improving the mechanical properties.

This is the reason of the strength and the elongation to failure variation by applying the hot rolling process. The strength and the elongation of the as-hot rolled ones are higher than that of the as-SPSed ones which can attribute to presence of the CNTs. Pores as process artifacts have a significant effect on the elongation to failure. Density is a proper factor in order to evaluate the process pores. With regard to the densities increase after hot rolling process for all samples, it is expected that the pores had minimum effect on the ductility.

In contrast, the as-rolled pure aluminum due to absence of the CNTs, depicted a larger grain size if compared to its as-SPSed counterpart. This grain size growth led to increase the ductility of the as-rolled sample with respect to as-SPSed one. It is attributed to the larger dislocation slip in the absence of the obstacles such the CNTs, the GBs, and pores. Moreover, it can be mainly due to the porosity reduction during the hot rolling process. Additionally, applying the pressure caused by the rolling simultaneously as well as high temperature resulted in more and effective rearrangement of the primary grains besides recrystallized grains. The incomplete dynamic recrystallization occurred in the as- hot rolled pure aluminum which leads to formation of new HAGBs and also LAGBs in different regions of the material. It seems evident that, grain boundaries ratio increases in as-hot rolled pure aluminum comparing to the as-SPSed counterpart. Additionally, grain refinement in the as-hot rolled pure aluminum is due to dynamic recrystallization which leads to formation of small new grain in the structure. As a conclusion, strength and elongation increase, simultaneously.

Note that, hindering effects of CNTs on dislocation slips typically leads to a remarkably decrease in elongation [1, 3]. In fact, the CNT nanoinclusions tend to situate at GBs, resulting to postpone or decrease the dynamic recovery of the dislocation. Consequently, suppress the GB migration/grain growth/sliding during materials synthesis. It is worth to note that by decreasing the aluminum grain size and subsequently the increasing of the GB fraction, the CNTs were predicated to be uniform dispersed in the hot rolled nanocomposites because of the presence of the CNTs around of the small grains. At the same time, in nanostructure materials by decreasing the grain size, lattice dislocation slip is limited so that the suppressing effects of Orowan looping of the generated dislocations at CNTs. Hindering the dislocation slip as well as the dislocation-controlled grain refinement process are affected by the presence of the CNTs.

Finally, hot rolling leads to the increase of the elastic modulus for all the materials as a direct consequence of the porosity reduction. Different mechanisms are presented and explained in literature in order to justify the strengthening effects of carbon nanotubes in different matrices. First of all, this is due to the general behavior of metal matrix composites, the load transfer from the metal to the nanotubes [27]. Grain refining due to the reinforcement addition causing grain boundary pinning results effective [28], strengthening of in situ formed carbide from the reaction between aluminum and damaged CNTs [18, 91], effect of Zenner due to the presence of the fine alumina particles [5, 23, 31, 92, 93], Orowan mechanism due to CNTs looping, and thermal mismatch between the matrix and the nanotubes are recognized as strengthening factors [29]. Obviously, these mechanisms are very effective if the nanotubes integrity is retained. Concerning the thermal mismatch strengthening, it is worth to note that there were mismatches between the predicted strength and the experimental values [13, 14]. Additionally, there is no countable evidence for thermal mismatch strengthening. Therefore, it is expected that the thermal mismatch strengthening does not give significant contribution in the high performance of the understudy composites. Moreover, once the nanotubes are strongly damaged during processing, the effectiveness of the mechanisms is reduced and the material strength drops. In general, load transfer and grain refinement are recognized as the main strengthening mechanisms in these composites [30, 31]. Again, especially the load transfer effectiveness is related to the nanotubes integrity. The latter is believed to be the main strengthening factor in this kind of materials contributing to the resistance increase for over 60% [24]. For the employed sintering temperatures, interfacial behavior between the nanotubes and the aluminum matrix is believed to act. This leads to a mixed mechanical-chemical bonding due to the formation of Al4C3nanoroads. The change from simple mechanical to mechanical-chemical bonding is believed to be beneficial for the strength increase. The optimal conditions were demonstrated to be obtained once a sintering temperature in the range 825–875 K is employed [32]. Here, the formation of Al4C3 at the interface leads to a volume expansion with consequent compressive stresses at the matrix-nanotube interface. This increases the interfacial friction, hindering the slippage of the carbon nanotubes in the matrix [26]. For these reinforcement percentages, agglomeration of nanotubes is negligible so the decrease of mechanical properties for the highest MWCNTs of the present study cannot be due to this aspect [33]. So, given the load transfer as the main strengthening mechanism, its efficiency increases as the carbon nanotube aspect ratio increases. So, once the nanotube damage is avoided, the strength increases for the mentioned mechanism [29, 34]. Once the reinforcement is damaged, the Orowan mechanism becomes predominant but it has lower strengthening efficiency with respect to the load transfer [35]. It can be also supposed that only a part of the nanotubes agglomerates leads to a decrease in the load transfer effect [36, 37, 38]. This can be due to the ultrasonication method allowing to obtain optimal conditions for reinforcing percentages lower than 1.

TEM observations indicate that during ball milling some of CNTs probably were bent and broken. Such finding previously was reported by Yoo et al. [94]. Therefore, CNT length is decreased in some cases. In this circumstance, the produced nanocomposites could strength via Orowan mechanism which is the result of the dislocation interaction and CNTs (those that have a length less than the critical length), broken Al2O3 nanoparticles, and Al4C3 nanoparticles. Additionally, the rolling process could assist to the uniform distribution of CNT, Al4C3 as well as Al2O3 nanoparticles. The strengthening contribution of Orowan mechanism in Al-CNT nanocomposites can be estimated from Eq. 5 [75, 95]:
$$ \Delta {\sigma}_{\mathrm{Oro}}=\frac{MGb}{2.36\pi }\ \left(\ln \left(\frac{\varnothing }{2b}\right).\frac{1}{\lambda -\varnothing}\right) $$
(5)
where ∅ is the diameter of the CNTs (30 nm), M is the Taylor factor (= 3.06 for Al [75]), G is the shear modulus (25.4 GPa for Al [96]), b is the Burgers vector (= 0.286 nm for Al [97]), and λ is the interparticle distance (\( \lambda =\frac{1}{2}\varnothing \sqrt{\frac{3\pi }{2{f}_{cnt}}} \)) [94] for sphere-shaped particles such as Al4C3 and Al2O3 and \( \lambda =\varnothing \sqrt{\frac{\pi }{2{f}_{cnt}}} \)) [94, 98] for rod-shaped reinforcements such as CNTs). At the same time, most of CNTs possess the length more than the critical length. The load transfer strengthening which derived from shear lag theory [99] follows the Kelly–Tyson model [78]. Thus, the stiffness of the MWCNTs is directly utilized. High aspect ratio of MWCNTs is favored with this model (aspect ratio for MWCNTs is about 100 [14]). Concerning the Al-MWCNT nanocomposites, the applied force generates the shear stress along the interface of CNTs and aluminum. The generated shear stress can be transmitted from the aluminum matrix to the CNTs. Therefore, it can conclude that the UTS strength of Al-MWCNT nanocomposite is strongly related to the CNTs length (l). The strength can be calculated by Kelly–Tyson (Eqs. 67) [45]:
$$ {\sigma}_{C- LT}={\sigma}_{\mathrm{CNT}}{f}_{\mathrm{CNT}}\left(1-\frac{l}{2\ {l}_c}\right)+{\sigma}_{\mathrm{Al}}^{\prime }{f}_{\mathrm{Al}}l>{l}_c $$
(6)
$$ {\sigma}_{C- LT}={\sigma}_{\mathrm{CNT}}{f}_{\mathrm{CNT}}\left(\frac{l}{2\ {l}_c}\right)+{\sigma}_{\mathrm{Al}}{f}_{\mathrm{Al}}l<{l}_c $$
(7)
where σC − LT is the strengthening contribution of the load transfer mechanism in tensile strength of the composite, σAl is the tensile strength of the matrix, \( {\sigma}_{Al}^{\prime } \) is the stress on the aluminum when the MWCNTs ultimately fail in the composite. σCNT is the strength of the CNTs (= 30 GPa), fCNT (0.5 wt. % ~0.52 vol. % and 1.5 wt. % ~1.56 vol. %), and fAl are the volume fraction of the CNTs and the aluminum respectively, and lc is the critical length, which can calculated by Eq. 8 [45]:
$$ {l}_c=\frac{d{\sigma}_{\mathrm{CNT}}\kern0.75em }{2{\tau}_{\mathrm{Al}}} $$
(8)
where τAl is the shear yield stress of the matrix (Kelly has been reported that τAl can be assumed the half of tensile yield stress of the aluminum matrix [78]), d is the average diameter of the CNTs (30 nm). The calculated σC − LT is the minimum tensile strength for the nanocomposites. In fact, in the denominator of Eq. 8, there is a coefficient that is related to the load transfer efficiency and is equal to 1 if the matrix can transfer the load to the fiber completely. Otherwise, it is below one due to weak interfacial bonding between aluminum and CNTs, or peculiar deformation behavior of the matrix. In this study, τAl is assumed 35 MPa, σCNT is 30 GPa. Therefore, the critical length of the MWCNTs is calculated about 12.85 μm. Thus, in current work, l > lc considering the value of l is 30 μm. In such condition, loading of the aluminum nanocomposites reinforced with CNTs involves maximum load transfer. Of course, it should be noted that most of the studied composites are expected to have a critical length larger than the CNT length [24, 75]. Based on the findings of Chen et al. [75], the strengthening behavior of the Al-CNT nanocomposites during plastic deformation contains three distinct regimes. First regimes was the so-called regime I where the CNTs has a small length or low aspect ratio (l < lc), usually smaller than 10. In this circumstance, the applied force is not transferred to the CNTs; consequently, the Al-MWCNT nanocomposite fractured because of the ductile failure of aluminum matrix. In fact, the very fine CNTs could act as obstacles in front of the movement route of the formed dislocations caused by the different coefficients of thermal expansion (CTE) at the interface between the Al/CNT and Al-Al4C3. It can be demonstrated that the Al-CNT nanocomposite was mainly strengthened by Orowan looping as well crystallite refinement. For CNT length larger than the critical length, l > lc, the load transfer strengthening mechanism has a significant role on increasing the strength. The strength predicted by Eq. 6 is very high. The contribution of the strengthening mechanisms involved in mechanical properties of the Al-CNT nanocomposites is given in Table 5. In our study, for as-rolled nanocomposites, the tensile strength was about 375 MPa and 440 MPa that is consistent with the predicted values by Eq. 6. On the other hand, the Orowan strengthening mechanism had a minimum effect on the strengthening of the nanocomposites. However, due to the presence of CNT with length smaller than the critical length that resulted in damaging or breaking of the CNTs during milling and rolling, the Orowan looping mechanism appears.
Table 5

The contribution of the strengthening mechanisms in the studied nanocomposites as a function of the CNT content

Material

CNT, vol.%

Calculated

Experimental

σGR, MPa

σC − LT, MPa

σC − Oro, MPa

σC , MPa

As-hot rolled pure Al

0

60 (26.6%)

165 (73.3%)

250

As-hot rolled Al-0.5 %CNT

0.52

7 (1.91%)

338 (92%)

21 (5.7%)

375

As-hot rolled Al-1.5 %CNT

1.56

11 (2.57%)

345 (80.79%)

71 (16.62%)

442

A large length or aspect ratio of CNTs leads to the CNT fracture mode during composite failure [80]. Therefore, the critical length (lc) depicts the minimum length required for a perfect load transfer. On the other hands, the used MWCNTs could be loaded up to their tensile strength values. When l ≥ lc, the strengthening mechanism of load transfer is dominant, while the Orowan strengthening mechanism has a secondary role in the nanocomposite strength. Considering to the Orowan strengthening mechanism origin from trap and looping the dislocation between hard nanosized particles; thus, it is difficult to pass through the long CNTs and consequently form the dislocation dinks around interface of long CNTs to aluminum. In such conditions, the dislocation loops caused by Orowan bowing are not produced. It is in agreement with findings [75]. Therefore, it is expected that the Orowan mechanism contribution could be much lower than load bearing mechanism contribution in the strengthening of Al-CNT nanocomposites. However, the presence of nanoparticles of alumina and Al4C3 which the last one is as a result of the damages entered into CNT structure leads to a Orowan contribution in strengthening of the nanocomposites, but it has lower strengthening efficiency with respect to the load transfer.

The fracture surfaces of the SPSed composites as a function of CNT content are shown in Fig. 16.
Fig. 16

Fracture surfaces of the SPSed a Al-0.5wt.% CNT and b Al-1.5wt.% CNT

A mixed brittle-ductile surface can be underlined. Cracks nucleate by microvoids between adjacent particles. These microvoids nucleate at the particle-reinforcement interfaces or at the particle triple junctions as a consequence of incomplete bonding of aluminum particle during SPS process [17, 21]. By comparing Fig. 16a, b, it could be worse by increasing the CNT content. Figure 17 reveals that a completely different fracture mechanism occurred in the hot rolled samples.
Fig. 17

Fracture surfaces of the hot rolled a Al-0.5wt.% CNT and b Al-1.5wt.% CNT

Figure 17a, b depicts that almost a shear ductile fracture in these two samples, which is characterized by shallow small elongated shear dimples oriented along the shear direction. In shear ductile fracture, slip bands impinge on the second phase particles and inclusions, causing local strain concentrations which nucleate the voids. The stress triaxiality impedes the growth of the formed voids, and due to the dominating shear stress state, the voids will undergo substantial shearing and the final failure is then caused by the internal void shearing mechanism [39]. Therefore, it can be said that the fracture mechanism changed to a shear ductile fracture due to hot rolling of the samples. In addition, the hot plastic deformation leads to the improving of the aluminum particle bonding. A more brittle facture occurred in the 1.5% MWCNT sample as a result of the higher amount of reinforcement [40]. The material is characterized by low amount of plastic deformation, dimples with decreased size and depth, and also more smooth surface without dimples formation.

4 Conclusions

In summary, nanocomposite sheets of Al-CNT were fabricated by a powder metallurgy route via the combination of spark plasma sintering and subsequent hot rolling process. By increasing CNT content, the obtained sheets depicted the best mechanical performance due to a tight interface and uniform dispersion of the CNTs into aluminum matrix. However, if CNT content exceeds 0.5, some agglomerates of CNTs along with a large number of crystallite of aluminum carbide formed. The hot rolling could assist the CNTs orientation, increasing the interfacial bonding between Al and CNTs, to obtain the uniformity of dispersion of the CNTs, and reaching nearly full density of the products. Additionally, a minor quantity of aluminum carbide, formed during the SPS and even after hot rolling process, could be very effective for achieving to a strong bonding between the Al and CNTs. The increase in the mechanical properties with a good balance of the strength and ductility were attributed to the effective and strong bonding of CNTs to the aluminum matrix. Moreover, it was underlined that the load transferring, Orowan, and grain size strengthening mechanisms are effective mechanisms on the increment of mechanical properties. The contribution of these mechanisms is varied by increasing the CNT content.

Notes

Compliance with ethical standards

Ethical statement

The ethical standards were respected. The paper is written and submitted following the rules of good scientific practice.

Conflict of interest

The authors declare they have no conflict of interest.

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© Springer Nature Switzerland AG 2019

Authors and Affiliations

  1. 1.Department of Material EngineeringIsfahan University of TechnologyIsfahanIran
  2. 2.State Key Laboratory of Metal Matrix CompositesShanghai Jiao Tong UniversityShanghaiChina
  3. 3.Department of Innovation EngineeringUniversity of SalentoLecceItaly
  4. 4.Educational workshop centerIsfahan University of TechnologyIsfahanIran
  5. 5.Institute of Materials and Machine MechanicsSlovak Academy of SciencesBratislavaSlovak Republic
  6. 6.Institute of Materials and Machine MechanicsBratislavaSlovakia
  7. 7.Department of PhotonicsFaculty of Physics University of KashanKashanIran

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