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Enhanced thermal conductivity of graphene nanoplatelet-polymer nanocomposites fabricated via supercritical fluid assisted in-situ exfoliation S. Mahdi Hamidinejad, Raymond K.M. Chu, Biao Zhao, Chul B Park, and Tobin Filleter ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b15170 • Publication Date (Web): 11 Dec 2017 Downloaded from http://pubs.acs.org on December 14, 2017

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ACS Applied Materials & Interfaces

Enhanced thermal conductivity of graphene nanoplatelet-polymer nanocomposites fabricated via supercritical fluid assisted in-situ exfoliation S. Mahdi Hamidinejad a, b, Raymond K.M. Chu a, Biao Zhao a, Chul B. Park a*, and Tobin Filleter *

b

a

Microcellular Plastics Manufacturing Laboratory, Department of Mechanical and Industrial Engineering, University of Toronto, Toronto, Canada M5S 3G8

b

Nano Mechanics and Materials Lab, Department of Mechanical and Industrial Engineering, University of Toronto, 5 King’s College Road, Toronto M5S 3G8, Canada

*

Corresponding Authors’ Information: E-mail: [email protected]; [email protected], Address: 5 King's College Road, Toronto, Ontario M5S 3G8, Canada

Abstract: As electronic devices become increasingly miniaturized, their thermal management becomes critical. Efficient heat dissipation guarantees their optimal performance and service life. Graphene nanoplatelets (GnPs) have excellent thermal properties that show promise for use in fabricating commercial polymer nanocomposites with high thermal conductivity. Herein an industriallyviable technique for manufacturing a new class of lightweight GnP-polymer nanocomposites with high thermal conductivity is presented. Using this method, high-density-polyethylene (HDPE)GnP nanocomposites with a microcellular structure are fabricated by melt mixing, which is followed by supercritical fluid (SCF)-treatment and injection molding foaming which adds an extra layer of design flexibility. Thus, the microstructure is tailored within the nanocomposites, and this improves their thermal conductivity. Therefore, the SCF-treated HDPE-17.6 vol.% GnP microcellular nanocomposites have a solid-phase thermal conductivity of 4.13±0.12 Wm-1K-1. This value far exceeds that of their regular injection-molded counterparts (2.09±0.03 Wm-1K-1) and those reported in the literature. This dramatic improvement results from an in-situ GnPs’ exfoliation and dispersion, and from an elevated level of random orientation and interconnectivity. Thus, this technique provides a novel approach to the development of 1 ACS Paragon Plus Environment

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microscopically-tailored structures for the production of lighter and more thermally conductive heat sinks for the next generations of miniaturized electronic devices. Keywords: Thermal conductivity, graphene nanoplatelets, polymer nanocomposites, supercritical fluid, microcellular structure

1.

Introduction

Heat dissipation functionality is extremely critical in high-energy density systems such as nextgeneration miniaturized electronic devices 1. The continuous development of the smaller, lighter, and faster electronic components of such devices means that the heat they generate needs to be efficiently dissipated by more compact and lightweight heat sinks. Lightweight, multifunctional, low cost, and highly thermally conductive polymer composites show promise for use as heat dissipation components 2. When compared with metallic and ceramic composites, polymer composites have an attractive array of properties, including ease of processing, superior resistance to chemicals and corrosion, and tailorable physical/mechanical properties3–5. The thermal conductivity of polymer composites is intensely affected by their interfacial thermal resistance and interfacial phonon scattering 6,7, by their dispersion and orientation, and by the type of fillers used 8. Conventionally, thermally conductive polymer composites are filled with a high loading (50−80 vol.%) of micro-size fillers to achieve target thermal conductivity values (>1 Wm-1K-1) 9. With such a high filler loading level, however, the amount of polymer matrix left to support the fillers and the composite’s structural integrity is insufficient. This leads to expensive and heavyweight composites, which are difficult to process. One promising way to address this drawback is to incorporate nanomaterials with extraordinary thermal conductivity, higher aspect ratios, and mechanical properties during the creation of these composites.

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With the recent advances in nanomaterials and their growing availability, the types and functions available for polymer composites have been significantly increased. This has increased the opportunities to develop polymer nanocomposites with superior thermal conductivity. Extraordinary heat transport properties in such nanomaterials as graphene, carbon nanotubes (CNT), and boron nitride nanotubes (BNNT) have driven the research of polymer nanocomposites. However, the expected dramatic enhancement of thermal conductivity by the incorporation of CNTs 8 and BNNTs 10 has not yet materialized in polymer nanocomposites, even at very high additive loading levels. Despite the excellent thermal conductivity reported for individual nanotubes11, CNTs and BNNTs have not been shown to substantially improve the thermal transport properties of polymer nanocomposites

8,10

. This has been attributed to the

nanotubes’ one-dimensional nature, which leads to their having anisotropic thermal conductivity in the axial direction

11–13

. However, it has been suggested

8,14,15

that 2D nanomaterials such as

graphene can be a more effective nanomaterial for polymer nanocomposites with a high thermal conductivity. In recent years, graphene has attracted a great deal of attention due to its exceptional mechanical, electrical, and thermal properties. Notably, the thermal conductivity of single-layer graphene has been reported, as ∼5000 W/ (m.K)

16–18

. However, the practical underpinning needed to

economically manufacture graphene-based polymer composites is missing. It has been extremely challenging to exploit graphene’s full potential. This has been due to the complexities that exist in the exfoliation, dispersion, and control of the GnPs’ orientation within the composites 19. Various strategies, such as in-situ polymerization alignment by electrical field

23

20,21

, GnP surface modification

, and the use of hybrid additives

2,9

9,22

, GnP

have all been proposed to

develop polymer composites with high thermal conductivity. Table 1 summarizes some of the recent advances made in the development of thermally conductive polymer nanocomposites. 3 ACS Paragon Plus Environment


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Nevertheless, all of these fabrication techniques have been batch-type processes. This makes them expensive, time-consuming, and not easily scalable. Furthermore, in most cases, the required additive loading levels remain rather high. On the other hand, SCF-treatment and physical foaming have shown promise in enhancing the electrically conductive polymer composites’ functionalities in different applications

3,24–31

.

Incorporating the optimum microcellular foaming structures into the conductive polymer composites can significantly reduce the product’s weight. At the same time, it can also add another degree of design flexibility to help control the polymer composites’ functional properties. During foaming, cell growth can change both the alignment and orientation fillers around the growing bubbles through the biaxial stretching of the polymer matrix

3,24,28,29,32,33

. Furthermore,

applying the SCF-treatment and physical foaming to the polymer composites can enhance the dispersion 31,32,34 and distribution 28,29 of the additives in the polymer matrix. It also can lower the fillers’ mechanical breakdown

29,30

during processing. In this way, the optimized SCF-treatment

and microcellular foaming can introduce tailored structures that support the various functionalities such

as

electromagnetic

interference

shielding

effectiveness

25,27,29,30,35,36

,

electrical

conductivity,3,24,28–30 and the dielectric properties of conductive polymer composites

3,24

.

However, to the best of our knowledge, no attention has been paid to the role of SCF-treatment in promoting heat dissipation in thermally conductive polymer composites. In contrast to the batch-type methods (Table 1), injection molding is a common and economically-viable industrial technology used to manufacture polymer parts. Therefore, injection molding combined with a SCF-treatment of polymer composites can be an easy solution to generate tailored microstructures that improve the heat dissipation properties in graphene-based polymer nanocomposites. However, to the best of our knowledge, no effort has yet been reported on the heat dissipation performance of injection-molded microcellular nanocomposites. 4 ACS Paragon Plus Environment

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Table 1. Thermal conductivity of various batch-type graphene/polymer nanocomposites materials

filler content

Epoxy/GnP

25 vol%

thermal conductivity [Wm-1k-1] 12.4

Epoxy/GnP

10 wt.%

1.53

Functionalization, solution mixing and a curing process

22

CBT/GnP

20 wt.%

7.1a

Solvent-free melting process followed by in-situ polymerization

37

PLA/hBN/GnP

16.65/16.65 vol%

2.77

Melt mixing and compression molding

2

PVDF/GnP

20 wt.%

0.562

High-shear solution mixing followed by bath-sonication

23

PA-6/GnP

10 wt.%

0.416

In situ polymerization with simultaneous thermal reduction

20

PA-6/GnP

12 wt.%

2.49

21

PC/GnP

20 wt.%

1.76

One-step in situ intercalation polymerization Melt mixing and compression molding

SBR/GnP

24 vol.%

0.48

Solution mixing and a sonication

39

PA6/hBN/GnP

GnP /1.5/20 wt.%

1.76

Liquid exfoliation, solution blending and hot-pressing

40

a

fabrication method Surface treatment + planetary centrifugal mixing

ref 9

38

In-plane Thermal conductivity

Our study demonstrates a SCF-assisted manufacturing method for producing thermally conductive GnP-HDPE nanocomposites by using injection molding to create heat dissipation components. The SCF-treated microcellular GnP-HDPE nanocomposites exhibited heat transport properties that were remarkably superior to those of the regular injection-molded nanocomposites 41,42

. Furthermore, they are comparable to the overall heat transport performances of bath-type

methods reported in the literature 2,9,20–23,37–39,41. This was due to the tailored microstructure of the GnP-HDPE nanocomposites created by the proposed technique.



2. 2.1.

Experimental Section Materials and sample preparation

A commercially available high-density-polypropylene (HDPE), Marlex® HHM 5502BN with a melt flow index 0.35 dg/min.-1 (230 ºC/2.16 kg) and a specific gravity of 0.955 g cm-3 (Group NanoXplore Inc. Montreal, QC, Canada) was used as the polymer matrix. The HDPE was filled with GnP grade heXo-g-V20, with average lateral dimensions of 50 µm, a surface area of 30 m2/g, and a specific gravity of 2.2 g.cm-3 (Group NanoXplore Inc. Montreal, QC, Canada). Commercial nitrogen (N2), supplied by Linde Gas, Canada, was used as the environmentally friendly SCF. The HDPE nanocomposites with a different GnP content were made by diluting the as-received HDPE-35 wt.% GnP masterbatch (Group NanoXplore Inc. Montreal, QC, Canada) with the asreceived neat HDPE through mixing in a twin-screw extruder (27 mm, L/D:40). A 50-ton Arburg Allrounder 270/320C injection molding machine (Lossburg, Germany), with a 30-mm diameter screw equipped with MuCell® technology (Trexel, Inc., Woburn, Massachusetts) was used to fabricate the GnP-HDPE nanocomposite samples. The mold contained a rectangular cavity with a fan gate after the sprue. The mold cavity dimensions were 132 × 108 × 3 mm. More details about the implemented mold in this study were reported by Lee et al. 43. Three different types of HDPE-GnP nanocomposites, namely injection-molded solid (IMS), injection-molded foam (IMF) and high-pressure-injection-molded foam (HPIMF), were prepared. The IMS samples were fabricated using the conventional injection molding process without the SCF-treatment and physical foaming. For the HPIMF and IMF samples, 0.4 wt% N2 (as the SCF), was injected into the barrel in its supercritical form using the MuCell module. The MuCell module is a built-on, commercially available, system for an injection molding machine to facilitate injecting the physical blowing agent into the barrel.


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In the IMF samples, after the GnP-polymer mixture was treated with the SCF, the mold cavity was partially filled with a gas-GnP-polymer mixture. In the HPIMF, the mold cavity was fully filled with the single-phase gas-GnP-polymer mixture. Then the filling step was followed by a composite melt packing step to re-dissolve the nucleated cells back into the melt. The nominal degrees of foaming in the IMF samples were controlled by partially filling the mold cavity. The processing parameters used in the injection molding of the IMS and IMF nanocomposites were optimized based on their microstructure integrity and thermal conductivity. Table 2 summarizes these processing parameters. A die cutter was used to cut disk-shape samples with a 20 mm diameter × 3 mm thickness from the injection-molded nanocomposites at a distance of 100 mm from the cavity gate. The schematic of the injection-molded parts has been presented in Figure S1. The IMF’s actual degrees of foaming were measured using samples via the waterdisplacement method (the ASTM D792-00) after fabrication. Table 2. Processing parameters used in injection molding of solid and foamed composites.

Parameter Melt temperature (°C)

IMS 210

HPIMF 210

IMF 210

Barrel pressure (MPa)

16

16

16

Screw speed (rpm)

300

300

300

12

12

12

Injection flow rate (cm s )

90

90

90

Mold temperature (°C)

75

75

75

Pack/hold pressure (MPa)

30

30

N/Aa

Pack/hold time (s)

15

60

N/A

Gas injection pressure (MPa)

N/A

24

24

N2 content (wt.%)

N/A

0.4

0.4

Degree of foaming (%)

N/A

N/A

7, 16, 26

Metering time (s) 3 -1

a

N/A: not applicable



2.2.

Characterization

The relative GnP powder’s defects, and an estimation of the number of layers it had, were determined using Raman spectroscopy (Renishaw, 532 nm laser excitation) (see Figure S2). Xray Photoelectron Spectroscopy (XPS) was also conducted on the GnP powder to identify its surface chemistry and functional groups, and also to measure the C/O ratio (see Figure S2). The results of Raman spectroscopy and XPS on the GnP powder are discussed in Supporting Information. An X-ray photoelectron spectrometer (ThermoFisher Scientific K-Alpha) equipped with an AlKα X-ray source was used to collect XPS data to analyze the qualitative defect density of the GnPs. To examine the exfoliation and dispersion of GnPs in the polymer matrix, Wide Angle X-ray Diffraction (WAXD) analyses were conducted on the injection-molded nanocomposites using a Rigaku MiniFlex 600 X-ray diffractometer (Cu Kα radiation, λ = 1.5405 Å). To further evaluate the level of exfoliation and dispersion of different samples, transmission electron microscope (TEM; FEI Tecnai 20) were conducted. The TEM samples were prepared by cryo-ultramicrotomy (Leica EM FCS). The microstructure and morphology of the fabricated samples were investigated using scanning electron microscopy (SEM; Quanta EFG250). The samples were frozen in liquid nitrogen, cryofractured, and sputter-coated prior to electron microscopy. The thermal conductivities of the GnP-polymer nanocomposites were measured using the transient hot disk method. A transient plane source (TPS) hot disk thermal constants analyzer (Therm Test Inc., TPS 2500, Sweden) was used to measure the samples’ thermal conductivity under ambient conditions with a Kapton (C7577) sensor. Measurements were taken based on the ISO/DIS 22007-2.2 standard. In this method, an electrically conductive double spiral disk-shape sensor made of nickel foil works as both a heater, to increase the temperature, and a dynamic thermometer to record the change in samples’ temperature as a function of time. The sensor is 8 ACS Paragon Plus Environment

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placed between two pieces of the sample and the increase in the samples’ temperature is evaluated by the analyzer to calculate the thermal conductivity (See Figure S3). Therefore, the generated heat will be dissipated in any direction (e.g. through-plane and in-plane) and the measured thermal conductivity is the overall (total) thermal conductivity 44,45.

3.

Results and discussion

3.1.

Microstructure and morphology of GnP-polymer nanocomposites

The IMS samples were fabricated without SCF-treatment and physical foaming. In the IMF samples, we first obtained a single-phase gas-GnP-polymer mixture, by SCF-treatment. When the mold cavity was partially filled with this mixture, physical foaming occurred due to the depressurization process. However, in the HPIMF samples the mold cavity was fully filled with the same mixture. Yet, as had happened with the IMF samples, the physical foaming occurred when the mixture had entered the mold cavity. However, the next step occurred under high pressure, and the nucleated cells re-dissolved back completely into the GnP-polymer mixture. The SCF-treatment and physical foaming produced a tailored microcellular structure, which increased the GnPs’ exfoliation and random orientation. A thinner skin layer also resulted. 3.1.1

Effect of SCF-treatment and physical foaming on GnP’s exfoliation and dispersion

To quantify the GnPs’ exfoliation level after the SCF-treatment, WAXD analyses were conducted. Figure 1 (a and b) shows the WAXD patterns for the neat HDPE, GnP powder, the IMS samples (HDPE-9 vol.% GnP) and their HPIMF and IMF counterparts. The diffraction peak at 2θ = 26.6° is characteristic of the (002) reflection of the graphite (I002), associated to the dspacing between the monolayer graphene sheets. By monitoring the (002) diffraction peak of the XRD pattern the stacking nature of GnP’s can be identified. As the ratio of exfoliated GnPs to stacked (unexfoliated) GnPs increases, the intensity of (002) diffraction decreases

46–52

. While a

low-angle shift of the (002) diffraction peak indicates GnP d-spacing expansion and intercalation 9 ACS Paragon Plus Environment


47,48

, the decrease in the intensity of (002) diffraction has been frequently used as the evidence of

exfoliation in literature 46–52. The SCF-treatment and physical foaming of the GnP-HDPE nanocomposites produced a 94% decrease in the intensity of the I(002) initial value’s diffraction peak, which corresponded to the untreated nanocomposites (IMS) (Figure 1c). This suggested very efficient exfoliation of GnPs which is in good agreement with the literature

46–52

. However, there was still a layered-GnP

structure retained in each flake 50, as evidenced by the presence of a small diffraction peak (I002), even after the SCF-treatment and physical foaming of the GnP-HDPE nanocomposites. Moreover, the SCF-treated GnP-HDPE nanocomposites’ (002) diffraction peaks shifted to somewhat lower angles, indicating a slight d-spacing expansion in the layered GnPs’ structure, based on Bragg’s law. To further support the GnPs exfoliation in the SCF-treatment technique we have also conducted WAXD on all of the samples over 2θ angles of up to 50° to examine the effect of GnPs orientation on the intensity of the (002) and (100) peaks. The intensity of the (002) and (100) peaks of layered structures such as GnP and hBN can be used to identify the orientation of these fillers within polymer composites

51,53,54

. Vertically and horizontally oriented flakes are

responsible for magnifying the (100) and (002) peaks respectively 53,54 as schematically shown in Figure 1b. The (100) peaks of all the samples are found to be very weak and they are not evident in Figure 1a. In a magnified XRD pattern over 2θ=40° to 50°, presented in Figure 1b, it was notable that the IMS and IMF samples had very small (100) diffraction peaks with similar intensity. However, the intensity of the (002) peak of IMS samples was more intense as compared to those of the IMF counterparts. This suggests that (i) the GnPs were horizontally oriented on the surfaces of IMS and IMF samples, and (ii) the decrease in the intensity of (002) peaks of the IMF samples is caused solely by exfoliation of GnPs and not by the orientation of GnPs 51,53,54. 10 ACS Paragon Plus Environment

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Figure 1c also shows that, once the GnP-HDPE nanocomposites had been treated with the SCF, the I002’s intensity considerably decreased. However, the degree of foaming (that is, the void fraction in percentage) did not significantly reduce the I002’s intensity in the range of 7-26 %. It is also noteworthy that the SCF-treatment provided almost the same level of exfoliation in the HPIMF samples as it had in the IMF samples. This was even after the nucleated bubbles had redissolved back into the composite melt under high pressure.



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Figure 1. (a) XRD spectra of neat HDPE, GnP powder, IMS samples (HDPE-9 vol.% GnP) and their HPIMF and IMF counterparts with various degrees of foaming; (b) magnified XRD pattern of Figure 1a over 2θ=40°-50° highlighted with light green, to examine (100) diffraction peaks and illustration of the GnPs’ orientation and their effect on the (002) and (100) diffraction peaks of the XRD pattern; (c) residual values (%) of I(002) (intensity of the (002) diffraction at 2θ = 26.5°) before and after SCF-treatment and physical foaming; (d) representative TEM micrographs of the IMS of HDPE-4.5vol.% GnP and; (e) IMF of HDPE-4.5vol.% GnP; (f) ideal conceptualization of various phenomenon resulting in further exfoliation and dispersion of GnPs in IMF samples. DF stands for degree of foaming.

In the IMF samples, the GnP-HDPE mixture is subjected to the SCF before being injected into the mold cavity. It is well known that SCF can help to enhance dissolution behavior

55

. Over a

sufficient duration, the SCF is capable of intercalating the graphitic layered structures. This


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weakens the nanoplatelets’ bonding force and makes their exfoliation easier. Moreover, the dissolution of the SCF in the HDPE melt creates a favorable interaction between the GnPs’ surfaces and the polymer melt which reduces system energy. This decreases interfacial tension between the GnPs and the polymer matrix, and it allows for a better GnP dispersion in the polymer melt

56

. Furthermore, the SCF’s plasticizing effect enhances the polymer molecules’

diffusivity. Furthermore, the SCF’s plasticizing effect enhances the polymer molecules’ diffusivity. This would likely increase the likelihood for polymer chains to penetrate the GnP nanoplatelets’ interlayer regions because of (i) the higher mobility of the chains by the plasticizing SCF dissolved in the polymer matrix; and (ii) the increase in the GnP nanoplatelets’ interlayer distances of the SCF-GnP intercalated structure. As a result, the GnPs’ exfoliation and layer separation is more effectively induced. To completely dissolve the SCF in the GnP-HDPE composite, it is necessary to maintain the GnP-HDPE/gas mixture’s single-phase throughout the injection molding process. This process was followed by a rapid depressurization to transform the dissolved and intercalatedSCF state into a gaseous state. During the phase transition, the expanding SCF can further separate and exfoliate graphene layers. Moreover, during the phase transformation, many small cells were generated between the platelets within the intercalated polymer/gas mixture. This led to further delamination and separation of individual platelets in the polymer matrix. Meanwhile, during the SCF’s depressurization and phase transformation, an additional driving force for the delamination and dispersion of the GnPs was generated. Nucleated cells growing near the GnPs acted like nucleating agents, and this further delaminated and uniformly dispersed the GnPs in the polymer matrix. Figures 2d and e, respectively show representative TEM micrographs of the IMS and IMF (containing 7% degree of foaming) samples. It is notable that agglomerated and thick GnPs in the IMS samples (Figure 1d) were further exfoliated to thinner layers after SCF-treatment 13 ACS Paragon Plus Environment


in the IMF samples (Figure 1e). This result is in a good agreement with the WAXD results, and provides further evidence of the higher level of exfoliation and better dispersion after SCFtreatment and physical foaming. Figure 1f shows the ideal conceptualization of the various phenomenon that resulted in further GnP exfoliation and dispersion in the IMF samples. Moreover, the dissolved gas in the composite melt reduced the melt’s viscosity and, therefore, lowered the shear stresses that were applied to the fillers. This helped to reduce the GnP’s mechanical breakdown. 3.1.2 Effects of SCF-treatment on the cellular microstructure, GnPs’ orientation, and skin layer Figure 2a shows the skin and core microstructure of the IMS HDPE-9 vol.% of GnP and that of its counterparts: HPIMF and IMF (7% degree of foaming). As expected, the IMS samples’ core and skin layers had completely solid structures. The IMF samples had a microcellular structure with a random cell morphology in both the skin and core layers, with an average cell size of 3 µm. The HPIMF samples’ structure was almost solid. And its cellular structure was barely visible due to the nucleated cells’ redissolution under high pressure. In IMF samples, cell growth caused different degrees of GnP rotation and displacement, which led to the GnP’s random orientation and further dispersion. To be specific, the GnPs get oriented more perpendicular to the radial direction with bubble growth and, consequently, the GnPs come to meet each other along the bubble surface. In other words, there was a greater chance of interconnectivity and direct GnP-GnP contact. This led to a particular morphology in which the IMF samples were greatly differentiated from the flow-induced structure found in the IMS samples. In the IMS samples’ skin layers (about 500 µm on each side), the GnPs were aligned in the machine direction (Figure 2a). This was due to the rapid cooling and the applied shear stresses in the direction of flow during the melt injection. This preferred filler alignment in the composites 14 ACS Paragon Plus Environment

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fabricated via injection molding has been well covered in the literature

24,28,29,57

. In the IMS

samples’ core layer, the GnPs followed the fountain flow orientation and were relatively randomly oriented. The HPIMF samples had the same skin-core morphology; however, the skin layer was thinner compared to their IMS counterparts’ (about 350 µm on each side). The composite melt’s lower viscosity, which was due to the SCF-treatment, had reduced the GnPs’ flow-induced orientation in the skin layer. This resulted in a lower skin layer thickness with oriented GnPs. Similar phenomenon has also been found in polymer/fiber composites

3,28,29,57

.

Conversely, in the IMF samples, the skin-core morphology and the oriented skin layer were hardly identified. This can be attributed to not only the composite melt’s lower viscosity but also to physical foaming. Figure 2b shows the ideal 2-D conceptualization of the evolution of the GnPs’ interconnectivity, orientation, and their further exfoliation due to SCF-treatment and physical foaming.



Figure 2. (a) SEM micrographs of skin and core regions for IMS, HPIMF and IMF HDPE-9 vol% GnP nanocomposites. Scale bars are all 10 µm; (b) ideal 2-D conceptualization of the evolution of GnPs interconnectivity, orientation and further exfoliation due to SCF-treatment and physical foaming; (c) SEM micrographs of IMF HDPE-9 vol% GnP nanocomposites showing different types of cells generated in the microstructure. FD stands for flow direction.

In the IMF samples, three different types of cells were found: (i) small cells that had nucleated in the polymer matrix, which led to cells with polymeric walls (shown in Figure 2c by green circles); (ii) cells that had nucleated at the edge, or on the surfaces, of the GnPs, which acted as nucleating 16 ACS Paragon Plus Environment

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agents, and which led to the cells being surrounded by a combination of polymeric walls and GnPs (shown in Figure 2c by yellow arrows); and (iii) cells formed by the phase transition of the SCF to a gaseous state, which led to cells encompassed only by GnPs (shown in Figure 2c by red arrows). To more clearly elucidate the microcellular structure, we have conducted additional electron microscopy imaging for IMF samples with lower magnifications which are presented in Figure S4. Because of the non-homogeneity of the structure with the dispersed and distributed GnP particles, the observed cells were quite non-homogeneous as shown in Figure 2c. This non-homogeneity cannot be explained with the non-homogeneous growth alone because some cell walls have a clean GnP surface. The bubbles must have been nucleated very non-homogeneously and nonuniformly. It is quite well accepted that the heterogeneous cell nucleation scheme will be preferred at the interface because of the lower activation energy for cell nucleation58. The clean surfaces of the cavities observed from Figure 2c indicate that those cells were nucleated at a surface of the GnP, based on the heterogeneous cell nucleation mechanism. It is clear from Figure 2c that most cells were nucleated this way. The size of this type of cells approximately ranges from 3µm to 20µm. But we could also observe the smaller cells (∼1µm) nucleated inside the polymer matrix alone because these cells were completely encapsulated by the polymer melt (see the green color circles in Figure 2c). On the other hand, Figure 2c also shows the other category cavities that were neither formed at the polymer-GnP interface, nor inside the polymer matrix. In fact, there are so many of these types of cavities that are observed from the SEM images. Since these cavities’ boundaries are GnP particles alone, not a polymer melt, these must have been formed by the expanding action of the SCF that diffused into the GnP layers before expansion (see cavities shown by red arrows in Figure 2c). The size of these bubbles is approximately ∼120µm. 17 ACS Paragon Plus Environment


3.2.

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Thermal conductivity

3.2.1. Effect of the GnP content on the thermal conductivity Figure 3a shows the total thermal conductivity of the IMS samples as a function of their GnP content. Their thermal conductivity is reported as a function of the final GnP content. The GnPs’ volume percent was calculated with respect to the total volume of foamed GnP-HDPE nanocomposites, including both gaseous and solid phases. As we had expected, in all of the samples, including the IMS, HPIMF and the IMF, the thermal conductivities of the GnP-HDPE nanocomposites were highly dependent on the GnP content. In the IMS samples, the total thermal conductivity corresponded to a 404% increase over the neat HDPE samples at an 18 vol.% of GnP and an increase of 21.3% per 1 vol.% GnP loading. This accorded with the enhancement efficiency of such traditional fillers as graphitic microparticles, which typically show an increase of ∼20% per 1 vol.% filler loading 9,59. However, this value for HPIMF and IMF samples was 31% and 46% respectively. On the other hand, for the IMF samples, introducing a 26% degree of foaming into the neat HDPE (i.e., at zero GnP loading), would decrease the thermal conductivity by 50% because of the creation of the voids with low thermal conductivity (0.026 Wm-1K-1 for ambient air)

60

. But,

interestingly, as the GnP loading increased, the detrimental effect of physical foaming on the IMF nanocomposites’ thermal conductivity became insignificant at around 7 vol.% GnP loading. The thermal conductivities of the IMF nanocomposites started to outpace those of the IMS and HPIMF composites at a GnP loading of more than 7 vol.%. We attributed this to a sufficiency of GnPs in the polymer nanocomposites to form thermally conductive paths. In other words, below a 7 vol.% GnP loading, the polymer matrix mediated between the GnPs. The result was a polymeric gap that broke the direct GnP-GnP contact. This caused the phonon scattering and high interfacial thermal resistance 9,61. 18 ACS Paragon Plus Environment

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3.2.2. Effect of GnP’s exfoliation and dispersion on the thermal conductivity It is interesting to note that the SCF-treated HPIMF counterparts’ total thermal conductivity was 563% greater than the neat HDPE samples at an 18 vol.% of GnP, and there was an increase of 31% per 1 vol.% of GnP loading. The increase in the HPIMF samples’ thermal conductivity over that of the IMS samples can be attributed largely to their higher level of GnP exfoliation, when compared to their IMS counterparts (Figure 1). In other words, at the same GnP loading level, the number of effective GnPs in the HPIMF samples was greater than the number of GnPs in the IMS samples due to a higher level of exfoliation with the SCF treatment. This increased the chance for direct GnP-GnP contact, which has a much lower interfacial thermal resistance than a polymer mediated structure would have (GnP-polymer contact) due to a lower amount of phonon scattering 9,61. Consequently, thermally conductive paths were formed more likely. It should be emphasized that the degree of foaming of HPIMF samples was almost negligible because of the high packing pressure (~300 MPa) used in the process. So, there would be negligible effect of the foaming on the thermal conductivity for the HPIMF samples.




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Figure 3. (a) The total thermal conductivity (λtotal) of IMS, HPIMF, and IMF GnP-HDPE nanocomposites as a function of the GnP content and; (b) the thermal conductivity of IMS, HPIMF, and IMF samples (HDPE-9 vol.% GnP) before (total) and after removing their skin (core); (c) the total thermal conductivity (λtotal) of IMS, HPIMF, and IMF GnP-HDPE nanocomposites as a function of the degree of foaming and the GnP content; (d) the total thermal conductivity (λtotal) of the samples as a function of the degree of foaming (GnP vol. % has been reported with respect to the polymer volume)



Page 22 of 34

3.2.3. Effects of GnPs’ re-orientation on the thermal conductivity of GnP-polymer nanocomposites It is also interesting to note that the IMF samples exhibited much higher thermal conductivities than their IMS and HPIMF counterparts, indicating that the local interconnectivity and the amount of direct GnP-GnP contact became much higher with foaming. This was due to the reduced orientation of the GnPs in the flow direction as well as the re-orientation of the GnPs surrounding the bubbles, caused by the foaming action that occurred in the IMF samples. This resulted in a lower interfacial thermal resistance than what would be found in a polymer mediated structure. For example, the total thermal conductivity of the IMF samples (with 7% degree of foaming) exhibited a higher increase of 46% per 1 vol.% GnP loading over the IMS samples with an increase of 20% per 1 vol.% GnP loading. Likewise, the total thermal conductivity of the IMF samples was also significantly higher than that of the skinless HPIMF samples. This outstanding improvement in the IMF samples’ total thermal conductivity over the IMS and HPIMF samples was attributed to the re-oriented GnPs’ microstructure in which the IMF samples were greatly differentiated from their IMS and HPIMF counterparts. Moreover, reduction of the GnPs’ orientation in the skin layer provided more isotropic heat transport functionality. The skin-core morphology, with highly oriented GnPs in the skin, was much more pronounced in the unfoamed IMS and HPIMF than in the foamed IMF samples (Figure 2). This resulted in a highly anisotropic heat dissipation property which deteriorated the product’s total thermal conductivity. It is worthy of noting Gong et al.’s claim

62

that too high

orientation of the conductive fibers will increase the percolation threshold even in the oriented direction. In fact, we observed an increased thermal conductivity from 1.20±0.01 to 1.40±0.04 Wm-1K-1 for the IMS samples with a 9 vol.% of GnP, after we removed, by machining, their skins with highly oriented GnPs (see Figure 3b). 22 ACS Paragon Plus Environment

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On the other hand, the higher thermal conductivity of the HPIMF samples over the IMS samples discussed in Section 2.3.2 may also have been affected by the lower thickness of the skin layer with highly oriented GnPs. Because of the reduced viscosity with the SCF treatment, the skin layer of the HPIMF would be reduced and, therefore, the thinner skin layer with highly oriented GnPs will increase the conductivity. As shown in Figure 3b, after removing the skins of the HPIMF samples (with 9 vol.% GnP), the thermal conductivity of the parts increased from 1.47±0.04 to 1.62±0.06 Wm-1K-1. However, the thermal conductivity of the IMF counterparts remained approximately constant after removing the skin layer (2.09±0.01 to 2.12±0.02Wm-1K-1). This is caused by the re-orientation of GnPs in the IMF samples due to the physical foaming leading to a more isotropic structure as compared to the IMS and HPIMF counterparts. In a nutshell, the enhanced thermal conductivity of the GnP-polymer nanocomposites with foaming was attributed to: (i) the reduced orientation of the GnPs in the flow direction; (ii) an increased local interconnectivity among the GnPs surrounding each bubble; and (iii) reduced orientation of the GnPs in a thinner skin-layer. It is notable that the crystallinities of the IMS, HPIMF and IMF samples are very similar (see Figure S5 showing Differential Scanning Calorimetry (DSC) and High-Pressure Differential Scanning Calorimetry (HPDSC)) on the HDPE-4.5 vol.% GnP). We also investigated the effects of the dissolved gas on the crystallinity of the HDPE-4.5 vol.% GnP using HPDSC. To investigate the non-isothermal crystallization in HPDSC, the HDPE-4.5 vol.% samples were heated and equilibrated at 200 °C for 30 min. The heating and thermal history removals were implemented under the N2 pressures of 1 and 48 bars. Then, the samples were cooled to 30 °C at a cooling rate of 10 °C/min, under N2 pressures of 1 and 48 bars in the HPDSC. We observed that the crystallinities and crystallization temperatures at different N2 pressures were very similar. This result was in good agreement with the DSC results. This can be attributed to the very fast 23 ACS Paragon Plus Environment


crystallization kinetics of HDPE which may not have been significantly affected by the parameters studied in this work. Therefore, we believe that the effect of crystallinity on the thermal conductivity in this study is negligible. 3.2.4. Optimal degree of foaming on the thermal conductivity Although foaming can enhance the thermal conductivity of the GnP-polymer nanocomposites, too high degree of foaming would be undesirable because of the non-conductive nature of the voids. This indicates that there exists an optimal degree of foaming to maximize the thermal conductivity. Figure 3 (c and d) shows the thermal conductivity variations with the GnP loading and the degree of foaming for the GnP-HDPE nanocomposites. When a 7% degree of foaming was introduced to the GnP-HDPE nanocomposites in the IMF samples, the total thermal conductivity was increased from 2.09±0.03 to 3.75±0.12 Wm-1K-1 at 17.6 vol.% of GnP. However, increasing the degree of foaming to beyond 7%, decreased the total thermal conductivity. This optimal behavior is attributed to the competing relationship between the favorable GnPs’ re-orientation effects (as discussed in Section 2.3.3) and the voids’ ultralow thermal conductivity. An excessive degree of foaming (that is, 16% and 26% in the current case) resulted in higher voids in the structure, which led to a lower thermal conductivity. The microstructures of the IMF samples with 7%, 16%, and 26% degrees of foaming appear to be similar to each other and there seem to be no particular features, according to Figure S6. Thus, we lean to conclude that the two competing mechanisms mentioned above govern the total thermal conductivity and that 7% was the optimal degree of foaming. 3.2.5. Solid phase thermal conductivity The total thermal conductivities of IMF samples were affected by two antagonistic parameters which included: (i) a constructive tailored morphological structure in the solid phase; and (ii) the negative impacts of insulating voids in the structure. To further analyze the net effects of the SCF24 ACS Paragon Plus Environment

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treatment and the foaming actions on the GnP-HDPE’s intrinsic thermal conductivity improvement, a theoretical model was employed in order to exclude the effect of voids and determine the thermal conductivity of the solid phase alone. This theoretical work was undertaken to quantitatively clarify the positive and negative effects of the voids on the total thermal conductivity. The convection that results from gas movement within the cells is negligible if the cell sizes are less than 4-5 mm 63. The contribution of radiation to the thermal conductivity of cellular plastics is less than 5% if their relative density are greater than 0.3

63

. It has also been reported that

carbonaceous materials as the infrared attenuated agents (IAAs) can block the radiation 64. The IAAs reported on in different studies can include surface-modified nano-graphite particulates 65, carbon nanotubes

66,67

, and dispersed graphene fillers

68

. Therefore, the contribution of the

radiative heat transfer does not apply to this study. In the IMF nanocomposite samples, the heat flux must pass through either the solid phase (GnPHDPE phase) or through the gaseous phase. Then, the total thermal conductivity (λtotal) of the IMF nanocomposites includes the solid conductivity (λsolid), and the gas conductivity (λgas) and is expressed as follows 60:

λtotal = λsolid + λgas

(1)

However, in confined spaces, the gas molecule collisions become lower. Thus, the gas conduction is governed by the energy transfer between the cell walls and the gas molecules. The Knudsen number (Kn) is defined to relate the dependency of the gas conductivity to the cell sizes as follows 69

:

Kn =

lmean d

(2)

where d is approximated by the cell size, and lmean is the mean free path of gas molecules, which is 68 nm in the ambient condition. The gas conductivity in polymeric foams is as follows 69: 25 ACS Paragon Plus Environment


Page 26 of 34

1 0 k gas (3) 1 + 2K n B where B is the energy transfer efficiency between the cell walls and the gas molecules,1.94. The k gas =

0 k gas is the bulk gas’s conductivity, which is 26 mWm-1K-1 for air 66.

We used the Maxwell-Eucken I model

70

in this study. This model is suitable for materials in

which the thermal conductivity of the dispersed phase is lower than the continuous phase (i.e., ksolid > kgas) such as polymeric foams 71. The Maxwell-Eucken I model is expressed as follows:

λtotal = k solid

2k solid + k gas − 2 (k solid − k gas )υ g 2k solid + k gas + ( k solid − k gas )υ g

(4)

where, ksolid and kgas respectively represent the thermal conductivity of the solid phase (GnPHDPE) and the gaseous phase. The ʋg is the degree of foaming (that is, the void fraction) of the IMF samples. Based on the cell sizes, the gas conductivities (kgas) of the IMF GnP-HDPE nanocomposites are calculated via Equation (3). The calculated kgas, the measured values of the ʋg, and the λtotal (that is, the total thermal conductivity of the IMF GnP-HDPE nanocomposites shown in Figure 3) are substituted in Equation (4). The thermal conductivity of the solid phase (ksolid) was then calculated and is plotted in Figure 4.


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Figure 4. Solid phase thermal conductivity (ksolid ) of IMS, and IMF GnP-HDPE nanocomposites as function of (a) the GnP content and; (b) the degree of foaming and the GnP content. DF stands for degree of foaming.

Figure 4a presents only the thermal conductivities of the solid phase, which were extracted from the IMF samples’ thermal conductivity using the Maxwell-Eucken I model. We note that the thermal conductivities of the solid phase in all of the IMF samples with various degrees of 27 ACS Paragon Plus Environment


foaming (i.e., 7%, 16% and 26%) coincided at approximately the same values. Thus, it was possible to evaluate the actual efficiency of the SCF-treatment and the physical foaming in the thermal conductivity enhancement of the GnP-HDPE nanocomposites. The thermal conductivity of IMF nanocomposites’ solid phase showed up to a 1,000% enhancement over the neat HDPE samples and an enhancement of 56% per 1 vol.% loading of GnP. In Figure 4b, the thermal conductivity of the IMF samples’ solid phase remained almost intact with a change in the degree of foaming. This occurred when the GnP-HDPE nanocomposites were first treated with the SCF, and then underwent the physical foaming. It is also noteworthy that the thermal conductivities of the IMF GnP-HDPE nanocomposites in the solid phase were higher than those of their IMS counterparts, even at GnP loadings of less than 7%. However, the difference between the thermal conductivities of the IMF’s solid phase and IMS samples was more pronounced with a higher GnP content.

4.

Conclusion

In our study, we introduced a new class of highly thermally conductive microcellular GnPpolymer nanocomposites. Microcellular nanocomposites containing highly exfoliated GnPs were developed by an industrially-viable technique of melt mixing followed by SCF-treatment and physical foaming in an injection molding process. This process provided a tailored structure that effectively supported the improved thermal conductivity of GnP-polymer nanocomposites. For example, the SCF-treated HDPE-17.6 vol.% GnP nanocomposites had a solid thermal conductivity of 4.13±0.12 Wm-1K-1 which was vastly superior to the values of their regular injection-molded counterparts (2.09±0.03 Wm-1K-1) as well as to those reported in the literature 41,42

. The reasons for this dramatic improvement includes the following: (i) a higher level of

GnPs’ exfoliation and dispersion in the polymer matrix; (ii) a decreased degree of GnP orientation from the reduced viscosity and the foaming action; (iii) an increased local interconnectivity 28 ACS Paragon Plus Environment

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among the GnPs surrounding each bubble; and (iv) a reduced skin-layer thickness. Our research shows that SCF-treatment of GnP-HDPE nanocomposites can add an extra layer of design flexibility in the manufacture of GnP-polymer composites with tailored morphologies and thermal conductivity. This design can be readily scaled up to an industrial level to make efficient and lightweight thermally conductive products for heat dissipation components in various miniaturized electronic devices.

5.

Associated content

Supporting Information: The schematic of the injection molded parts, defect density of GnPs, Raman spectroscopy of the GnPs, XPS measurement of the GnPs, the schematic of the setup for measuring the thermal conductivity, SEM micrographs of the FIM samples, DSC and HPDSC of the IMS, HPIMF and IMF samples.

6.

Acknowledgment

The authors gratefully acknowledge Group NanoXplore Inc.’s donation of GnP powder and GnPHDPE masterbatch, and the financial support of the Natural Sciences and Engineering Research Council of Canada (NSERC).

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Abstract:

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