CNTs nanocomposites by addition of mechanically activated graphite

Composites: Part A 68 (2015) 177–183

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Improvement of mechanical properties in aluminum/CNTs nanocomposites by addition of mechanically activated graphite S. Khorasani a,⇑, S. Heshmati-Manesh a,b, H. Abdizadeh a,b a b

School of Metallurgy and Materials Engineering, University of Tehran, Tehran, Iran Center of Excellence for High Performance Materials, University of Tehran, Iran

a r t i c l e

i n f o

Article history: Received 4 April 2014 Received in revised form 28 September 2014 Accepted 1 October 2014 Available online 17 October 2014 Keywords: A. Metal–matrix composites (MMCs) B. Adhesion B. Physical properties D. Mechanical testing

a b s t r a c t Carbon nanotubes reinforced aluminum nanocomposite was prepared by ball milling route. CNTs were initially mixed with mechanically amorphized graphite. Specimens were analyzed by X-ray diffractometry and Raman spectroscopy. Crystallite size and dislocation density were calculated by modified Warren–Averbach method. Carbide formation was semi-quantitatively investigated via Raman spectroscopy. A band located in 950 cm1 was considered to be corresponded to Al4C3. Hardness of the samples was also evaluated using a Vickers micro-hardness tester. The hardness strengthening contributions were modeled to evaluate interfacial bonding between CNTs and the aluminum matrix. In specimens, including amorphized graphite, hardening was due to both work hardening and second phase strengthening otherwise, only due to work hardening. It was deducted that the amorphized graphite has a major role for mechanical properties improvement. This seems to be due to the formation of aluminum carbide at the interface which consequently increases adhesion of CNTs to aluminum. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction Since the discovery of carbon nanotubes (CNTs) in 1991, they have been considered as promising reinforcements for nanocomposites, due to their exceptional mechanical and physical properties [1–3]. Since metallic materials have good wear resistance and high-temperature stability as compared to polymeric materials, metal matrix composites with CNT reinforcements seem to find wide field of applications as (multi-functional) structural materials [4]. Among many candidate matrix materials for lightweight highstrength composites, aluminum has been considered preferentially because of its relatively low density and reasonable mechanical properties. The dispersion and wettability problems are the main obstacles for good metal–CNTs composites due to the strong Van der Waals forces of attraction between these long and thin tubes [5]. Attempts at producing Al–CNTs by melt stirring were not satisfactory as a consequence of poor Al–CNTs wettability, high viscosity of Al and strong agglomeration of CNTs [4]. Current approach for dispersing CNTs in metal matrices has been reported

⇑ Corresponding author. Tel.: +98 9360171502, Postal code: 14199-93745. E-mail addresses: [email protected] (S. Khorasani), [email protected] (S. Heshmati-Manesh), [email protected] (H. Abdizadeh). http://dx.doi.org/10.1016/j.compositesa.2014.10.009 1359-835X/Ó 2014 Elsevier Ltd. All rights reserved.

to be conventional mechanical mixing techniques. A number of research groups, investigated the use of ball milling as a mechanical dispersion technique [6–8]. Electron microscope images showing well dispersed CNTs, proved the process to be promising [8–10]. However, concerns about the possible damage and/or amorphization of the CNTs under the harsh milling conditions have been raised [9,11], and therefore, optimization of milling conditions is necessary. Current authors evaluated different milling conditions for the case of aluminum in order to optimize milling process [12]. Although, only limited works were found to focus on wettability and interface phenomena between CNTs and aluminum [13,14]; it is well-understood that formation of carbide layer in the interface can significantly improve the interface strength and wettability [15–18]. Several researchers have utilized high temperatures consolidation processes to promote bonding between the aluminum matrices and carbon [6,11,19,20]. As a consequence, the processes tend to raise expenses. Moreover, complete reaction of CNTs with aluminum at elevated temperatures is inevitable [21–24]. The aim of the present study is to overcome the wettability problem and simultaneously to avoid and minimize destruction of CNTs. For this purpose, a novel technique was gained in which mechanically amorphized graphite was added to the composites to promote carbide formation at low temperatures.

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2. Experimental procedure

3. Results and discussion

Multi-walled carbon nanotubes (average dimensions: 30–50 nm outer diameter, 5–12 nm internal diameter, and 10–20 lm in length) were first subjected to a mixing process with previously milled graphite by conventional planetary ball mill with a ball to powder ratio of 5:1 and a rotation speed of 250 rpm for 5 h. The weight ratio of graphite to MWCNTs was 1:6. 2 or 3 drops of ethanol was used to avoid agglomeration. Milled graphite was used as an activated carbon phase in order to encourage carbide formation, and the mixing process was also carried out by milling in mild conditions (mentioned hereinbefore) such that the activated carbon phase could be distributed uniformly between CNTs, and even some lattice defects to be induced to CNTs’ structure. The lattice defects can further act as preferable sites for carbide phase nucleation. The batch of CNTs which were mechanically mixed with amorphized graphite will be recalled hereinafter as processed CNTs. Graphite had been milled at harsh conditions to produce amorphous and activated carbon phase. A ball to powder ratio of 15:1 with 5 mm diameter balls, a rotation speed of 300 rpm and time duration of 15 h were employed. Different mixtures of commercially pure aluminum (+170 mesh) with primary multi-walled carbon nanotubes and processed carbon nanotubes were provided as summarized in Table 1. The mixtures were milled in a high energy planetary ball mill with hardened chromium steel media and balls (10 mm diameter) with a rotation speed of 250 rpm for 5, 10 and 20 h. The milling operation was carried out under argon atmosphere and with a constant ball to powder ratio of 10:1. 0.5 ml of ethanol was added as a process control agent (PCA), which had been already indicated by the authors to act as an effective PCA [12]. For diagnosing the work-hardening effects, another three batches of pure Al powders were ball milled under the same conditions. The milled mixtures were cold pressed in a compaction die with an internal diameter of 12.5 mm at 650 MPa and then followed by sintering at 500 °C for 2 h. The structural changes of sintered samples during the process were determined by X-ray diffraction (XRD). A PANalytical X’Pert PRO diffractometer with Cu Ka1 radiation (k = 0.15406 nm) was used for the measurements. The crystallite size and dislocation density were determined by the modified Warren–Averbach method which has been described in more details elsewhere [12]. For structural analyzing of carbon phases as well as carbide investigation on sintered samples, Raman spectroscopy was accomplished. All Raman spectra were recorded at room temperature with an excitation wavelength laser of 785 nm, focused on the sample to a spot size of 2 lm diameter. Distribution of different phases in powder mixtures and processed CNTs were evaluated by scanning electron microscopy (SEM-AIS2100, Seron Co.) under an acceleration voltage of 24 kV, and field emission scanning electron microscopy (FE-SEM, Hitachi S4160), respectively. The mechanical response of the sintered specimens was also characterized using a Vickers micro-hardness tester (Load = 100 g and Dwell time = 10 s).

3.1. Processing on nanotubes

Table 1 Nomenclatures of the specimens. Mixtures

15.0 g Al 14.7 g Al + 0.3 g CNT 14.95 g Al + 0.05 g milled graphite 14.65 g Al + (0.3 g CNT + 0.05 milled graphite)a a

The processed CNTs.

Milling time duration 5h

10 h

20 h

5-Al 5-CNT 5-mg 5mCNT

10-Al 10-CNT 10-mg 10mCNT

20-Al 20-CNT 20-mg 20mCNT

It is well-known that CNTs have poor adhesion to metallic phases. However, with formation of Al4C3 at the interface, the contact angle would be reduced from P135° to 45° [15–17]. Ci et al. reported formation of Al4C3 after sputtering of aluminum on CNTs followed by annealing in the temperature range of 723–1223 K [13]. Formation of carbide in the interface usually needs high temperature processing [4]. However, fully conversion of CNTs to Al4C3 is highly probable in these circumstances [21,22]. Therefore, a new approach was needed to be designed in order to avoid high temperature requirement. Severely milled graphite was used as highly activated (amorphized) carbon to encourage carbide formation during low temperature processing. In Fig. 1a, Raman spectra of graphite are shown before and after the milling. Obviously, extreme reduction in peak intensities is due to amorphization. According to Table 2, the ratio of ID/IG which has been reported as a criterion to lattice defects density [25], is increased significantly after milling of graphite. The milled graphite was then mechanically mixed accompanied by as-received carbon nanotubes. The mixing process was carried out by low energy planetary ball mill for a short time of 5 h. The mixing processes were accomplished mechanically so that some lattice defects could be also induced to CNTs structure and subsequently act as potentially preferred sites for carbide nucleation [13]. In Fig. 1b Raman spectra of CNTs are depicted before and after the mixing process. No significant change was observed in spectra of asreceived and processed CNTs, as well as in ID/IG ratio (Table 2). Therefore, one can assume the equality of defect density in both as-received CNTs and the processed CNTs for further compositing process. SEM investigations revealed a good distribution of CNTs when mixed with amorphized graphite (Fig. 2a). Even though some agglomerated graphite were observed. Most of the processed CNTs were broken and reduced in their length (Fig. 2a). The SEM image of the as-received carbon nanotubes in Fig. 2b clearly shows their longer length when compared with the processed CNTs in Fig. 2a. 3.2. SEM images of the milled powder mixtures Distribution of CNTs between powder particles was studied by SEM images. In accordance with Fig. 3, no trace of CNTs agglomeration was observed and only few non-clear CNTs can be distinguished in 5-CNT and 5-mCNT samples, while no CNTs could be recognized among other mixtures milled for longer period of time. Liao et al. stated CNTs with short length can easily embed in aluminum powders [25]. Also, Wang et al. mentioned a mechanism for thread-like carbon nanotubes to be buried beneath aluminum powder surface [7], which as a consequence yields to a better distribution of CNTs within aluminum matrix in both cases [7,25]. As seen in Fig. 3a and b, buried (yellow arrows) and embedded (dark blue arrows) CNTs can be distinguished respectively. Further milling would gradually result in complete burying and/or embedding of CNTs into soft aluminum matrix (Fig. 3c–f) which could be the reason why CNTs could not be recognized through other samples. Even though Wang et al. reported existence of buried CNTs after 48 h of milling [7], burying of CNTs in this study occurred in shorter period of time because of optimized milling conditions [12]. 3.3. Raman spectra of composites Since aluminum metal does not have any Raman active modes, thus any other phases, even in very small amount, can be clearly

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Fig. 1. Raman spectra of (a) graphite and (b) CNTs, before and after milling. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Table 2 ID/IG of graphite, before and after milling.

Graphite Carbon nanotubes

Before

After

0.654 2.126

1.111 2.058

marked by Raman spectroscopy which has Raman active modes, such as carbon or carbide. Fig. 4 shows Raman spectrum of the as-milled 20-CNT. The ratio of ID/IG is markedly reduced after 20 h of milling with aluminum. Additionally, some new bands can be observed in the range of 200–1000 cm1, which is believed to be attributed to the carbide nucleation. In this case, carbide might have nucleated on CNTs defects, thereby resulting in reduction of lattice defects density (ID/IG). Although limited investigations on Raman spectrum of aluminum carbide are available; Sun et al. calculated Raman spectrum of Al4C3 with ab initio method. Their calculations result in 6 Raman active modes (3Eg + 3A1g). Raman shifts for Eg mode were determined as 247.89 cm1, 282.98 cm1 and 709.16 cm1, and for A1g mode as 337.95 cm1, 481.65 cm1 and 844.70 cm1 [26]. Raman spectrum of the as-received pure aluminum was also compared for better understanding of arisen new bands in Fig. 4. Amongst the new bands in 20-CNT’s Raman spectrum, a strong peak at 950 cm1 is observed which is absent in the as-received pure aluminum spectrum. This band was considered to be attributed to the Al–C bond. Therefore, A1g must have been blue-shifted 100 cm1, in comparison with calculated spectrum. Even though the calculated spectrum by Sun et al. was in good agreement with their experimental data, but it should be noted that they had used a laser with 514 nm excitation wavelength while the whole Raman spectra in this study were obtained by 785 nm excitation wavelength. Observation of such a shift by changing in excitation wavelength has been already reported in the literature [27–29]. Another reason for blue-shifting could be owing to distortions and narrow

thickness of formed carbide [30–33] (Reader is referred to Ref. [14] for more details about thickness of the formed carbide at the interface of CNTs). The milled samples were compacted and sintered at a low temperature of 500 °C for 2 h were provided according to Table 1. Table 3 semi-quantitatively analyzes carbide formation on the basis of normalized Raman spectra recorded from sintered specimens [34]. Since the desirable formed carbide is the one which nucleates at the interface of CNTs with aluminum matrix, the peaks in Raman spectra were normalized in respect to G band which is owing to the layered structures of graphite and multi-walled carbon nanotubes. The formed carbide is appreciable in 10-mCNT, 20-mCNT, 10-mg and 20-mg wherein amorphized graphite was added and milled for more than 10 h. A similar observation was reported by Arik [35,36]. Although the primary defect density of both specimens containing as-received and the processed CNTs, prior to mixing with aluminum, were equal (Fig. 1b and Table 2), and have the same number of potential sites for carbide nucleation (regardless to the added graphite), but formed Al4C3 was more appreciable for specimens containing the processed CNTs rather than as-received CNTs. This matter can illuminate the effect of amorphized graphite added to the processed CNTs.

3.4. X-ray diffractions of sintered composites XRD patterns of the sintered samples are indicated in Fig. 5. The peaks in XRD patterns were characterized in accordance with Joint Committee on Powder Diffraction Standards (JCPDS). Only pure aluminum (JCPDS No. 00-004-0787) was detected in sintered samples (Fig. 5a–c), except for 10-mCNT and 20-mCNT (Fig. 5d). For the case of 10-mCNT, both Al4C3 (JCPDS No. 35-0799) and c-Al2O3 (JCPDS No. 00-010-0425) have been crucially formed (Fig. 5d). While for the case of 20-mCNT only two new peaks have been appeared, which may belong to either Al4C3 or c-Al2O3. But, according to Table 3 in which significant amount of carbide was

Fig. 2. SEM image of (a) the processed CNTs (yellow arrows indicate agglomerated graphite) and (b) the as-received CNTs. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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Fig. 3. SEM images of the as-milled (a) 5-CNT, (b) 5-mCNT, (c) 10-CNT, (d) 10-mCNT, (e) 20-CNT and (f) 20-mCNT samples. The yellow and the dark blue arrows show buried and embedded CNTs, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Table 3 Semi-quantitative analysis of carbide in respect to CNTs and graphite. Nomenclature

I950/IG

5-mCNT 10-mCNT 20-mCNT 5-CNT 10-CNT 20-CNT 5-mg 10-mg 20-mg

0.156 0.333 0.520 0.200 0.216 0.197 0.200 0.301 0.328

Fig. 4. Raman spectrum of the as-milled 20-CNT sample which is compared with the as-received CNTs and pure Al. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

estimated for sintered 20-mCNT, the two new peaks are expecting to be related to Al4C3. 3.5. Mechanical tests Various mechanisms that contribute to the plastic deformation resistance can be determined by superposing as the following equation [19,37,38]:

H ¼ HPN þ HSS þ HC þ HD þ HP

ð1Þ

where H is the measured hardness, HPN is the Peierls–Nabarro hardness contribution, HSS the solid solution contribution, HC the hardness contribution caused by crystallite size, HD the forest dislocation contribution and HP the hardness contribution related to reinforced particles. For the case of CNTs or graphite as reinforced particles, HP will be appeared providing that interface boding between CNTs and aluminum is strong enough [19]. Since HPN and HSS (in spite of other terms) take action on the atomic scale, can be assumed to be smooth on the scale of the coarser mechanisms [19,37,38]. Therefore, the Eq. (1) can be expressed as the following equation:

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S. Khorasani et al. / Composites: Part A 68 (2015) 177–183 Table 5 Estimation of constants in Eq. (5) by the least square method. HL (HV)

k1 (nm HV) 7

4.641

9.954  10

k (m HV) 2.107  10+3

Fig. 5. XRD patterns of the sintered samples of (a) pure Al, (b) mixtures of Al with as-received CNTs, (c) mixtures of Al with amorphized graphite and (d) mixtures of Al with processed CNTs.

Table 4 Results from Warren–Averbach calculations with the average microhardness (Ave.) and standard deviation (S.D.). Nomenclature

Al (un-milled)a 5-Ala 10-Ala 20-Ala 5-CNT 10-CNT 20-CNT 5-mg 10-mg 20-mg 5-mCNT 10-mCNT 20-mCNT a b

q (1014 m2)

0 5.99 8.56 10.12 3.52 25.00 7.19 6.80 6.34 6.02 4.95 9.83 8.59

d (nm)

85.27b 46.84 44.21 41.01 56.85 35.33 45.02 64.95 44.43 44.14 65.97 49.71 41.90

Hardness (HV) Ave.

S.D.

28.05 79.91 80.50 83.95 58.17 116.96 76.73 58.55 82.21 95.12 57.22 165.31 126.38

1.13 9.08 1.68 1.00 3.09 4.53 1.82 3.68 1.73 3.26 4.02 4.94 6.74

Used for estimating HL, k1 and k. Calculated by Scherrer equation.

H ¼ HL þ HC þ H D þ H P

ð2Þ

where HL is the hardness contribution due to lattice resistance. The hardness contributions by crystallite size [39–41] and forest dislocations [19,42] is described as follows, respectively: 1

ð3Þ

1=2

ð4Þ

H C ¼ k1 d H D ¼ kq

where k1 and k are constants, d is the crystallite size and q is the dislocation density. The crystallite size and dislocation density were calculated by the modified Warren–Averbach method. The values of

Fig. 6. Contribution of each term in Eq. (2) for the composite samples milled for (a) 5 h, (b) 10 h and (c) 20 h. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

the mean crystallite size, dislocation density and hardness for consolidated samples are determined in Table 4. By substituting Eqs. (3) and (4) in Eq. (2), the strengthening contribution model for the first four specimens mentioned in Table 4 would be written as: 1

H ¼ H L þ k1 d

þ kq1=2

ð5Þ

Since these samples are made of pure aluminum, the last term related to the second phase has been eliminated. Having considered Eq. (5) and the data of the first four sample in Table 4, three constants of HL, k1 and k were estimated by the least square method (Table 5). Detail constitutive equations are given in Appendix. The estimated HL is in good agreement with Peierls–Nabarro stress for a/2 h1 1 0i edge dislocation in {1 1 1} plane of aluminum [43]. Fig. 6 illustrates contribution of hardness for each term in accordance with Eq. (2) and Table 4. The contribution of second phase strengthening is observed only for specimens in which amorphized graphite was added and milled for more than 10 h (Fig. 6b and c). Lack of HP for specimens milled for 5 h (Fig. 6a) and those containing amorphized graphite, indicates that the added second phase (either CNTs or the milled graphite) behaves as PCA [19]; that is

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due to poor interface bonding between the second phase and the matrix. It should be noted, in samples which revealed HP, appreciable formed carbide was estimated by Raman spectroscopy (Table 3). Any particle which has strong enough interface with matrix, would give rise to reinforcement strengthening [19]. One should be noticed that, although there is a possibility for formation of carbide as an isolated particle through aluminum matrix, but their effect as a second reinforcement particle can be eliminated due to the small amount of added graphite in comparison with CNTs. Nevertheless, HP measured for 10-mg and 20-mg (in which only graphite has been added) are less than those measured for 10-mCNT and 20-mCNT, as seen in Fig. 6b and c. In Fig. 6b, the higher measured hardness for 10-mCNT against 20mCNT may be owing to the presence of oxides in 10-mCNT which can act as hardener participates simultaneously with CNTs [44,45]. It should be noted that increasing temperature or subsequent processes (such as hot rolling and hot extrusion) during consolidation in other studies were a feasible option to yield increasing mechanical properties of composites by bonding between CNTs and aluminum [6,20,21,46]. However, full conversion of CNTs to carbide and destruction of tubular structure has been widely reported in these cases [21–24]. In this study, although the sintering temperature was kept at relatively low level and no other subsequent process was employed, second phase strengthening in Al–CNT composites was successfully achieved due to addition of activated graphite. 4. Conclusion Carbon nanotubes reinforced aluminum nanocomposites were prepared by ball milling. Carbon nanotubes were distinctively subjected to a mechanical mixing process with amorphized graphite. Mechanical mixing process caused breaking of CNTs and reduction of their length. The length reduction of CNTs gave rise to embedding of CNTs within the soft aluminum matrix and subsequently their better distribution. Raman spectroscopy of samples revealed formation of considerable amount of carbide phase in specimens which contained amorphized graphite and were milled for more than 10 h. Microhardness testing of the same samples showed good adhesion between the processed CNTs and the aluminum matrix. It is believed that the adhesion could be due to the carbide formation in the interface. To the best of our knowledge, this is the first time that CNTs adhesive to a metal matrix has been improved by addition of activated carbon, and without employing high consolidation temperatures. Appendix A In order to determine the constant coefficients of a linear function, as z = ax + by + c, with a given data set (x1, y1, z1), (x2, y2, z2), . . ., (xn, yn, zn), where n P 3, function P should have the least square error, i.e.:

P¼

n X 2 ½zi  ða þ bxi þ cyi Þ ¼ min

ðA1Þ

i¼1

To obtain the least square error, the unknown coefficients a, b and c must yield zero first derivatives.

8 n X @P > > ¼ 2 ½zi  ða þ bxi þ cyi Þ ¼ 0 > > > @a > i¼1 > > > n < @P X ¼ 2 xi ½zi  ða þ bxi þ cyi Þ ¼ 0 @b > > i¼1 > > > n > X > @ P > > ¼ 2 yi ½zi  ða þ bxi þ cyi Þ ¼ 0 : @c i¼1

ðA2Þ

By expanding the above equations, we will have:

8X n n n n X X X > > z i ¼ a 1 þ b xi þ c y i > > > > i¼1 i¼1 i¼1 i¼1 > > > n n n n > i¼1 i¼1 i¼1 i¼1 > > > n n n n > X X X X > > > yi zi ¼ a yi þ b xi yi þ y2i : i¼1

i¼1

i¼1

ðA3Þ

i¼1

The unknown coefficients a, b, and c can be obtained by solving the above linear equations.

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