Journal of Polymers and the Environment

, Volume 26, Issue 9, pp 3574–3589 | Cite as

Thermally Stable Pyrolytic Biocarbon as an Effective and Sustainable Reinforcing Filler for Polyamide Bio-composites Fabrication

  • Emmanuel O. Ogunsona
  • Amandine Codou
  • Manjusri Misra
  • Amar K. Mohanty
Original Paper


Natural fibers are limited in their use as reinforcement to commodity polymers. They cannot be used to reinforce engineering polymers due to their low thermal stability at high processing temperatures. This study presents an approach to successfully reinforce polyamides using a derivative of natural fibers as reinforcement without the effects of thermal degradation during melt processing. Biocarbon from miscanthus fibers was used to reinforce polyamide 6 up to 40 wt%. At 40 wt% filler content, the tensile and flexural strengths increased by 19.6 and 47% respectively in comparison to the neat polyamide. The moduli were also increased by 31.5 and 63.7% respectively. A maximum increase in impact strength of 43.7% was achieved at 20 wt% biocarbon loading. The morphology of the tensile fractured samples showed stretched polyamide ligaments attached to the biocarbon particles, indicating the presence of interaction between filler and matrix. Interestingly, more bonded interfaces were observed between the polyamide and biocarbon particles with increasing biocarbon content possibly stemming from increased biocarbon surfaces with functional groups. These composites show great potential to substitute in part or whole, some particulate filled polyamides currently used in the automotive industry.


Polyamide Biocarbon Polymer composites Reinforcements Bio-composites 


The new regulations in the automotive industry such as the Corporate Average Fuel Economy (CAFE) standard, stress a new driving force for all the automakers to improve on fuel efficiency which has resulted in the reduction in weight of their vehicles. For this, the replacement of all the non-renewable and high density fillers such as talc, glass and carbon fibres by sustainable and lower density fillers such as natural fibres are getting more attention [1]. Indeed, the effective use of bio-based fillers as sustainable and renewable alternatives for the fabrication of polymer composites has been receiving great importance in recent years due to the concerns about environmental and diminishing petroleum resource related issues. Natural fibers, wood flour and cellulose are some of such fillers which have been well explored, engineered and commercialized in different applications [2, 3, 4, 5]. However, most of these have been used as reinforcement in commodity plastics such as polypropylene (PP), polyethylene (PE), polystyrene (PS) and poly lactic acid (PLA). The use of these bio-based fillers in engineering plastics has been problematic because of their low thermal stability and hygroscopicity. Natural fibers begin to degrade around 180–200 °C and cannot be successfully incorporated into engineering plastics without thermal degradation. They easily absorb moisture from the surrounding environment and can be detrimental to moisture sensitive plastics like polyamide and PET during melt compounding at higher temperatures. Polyamide, with a processing temperature range of 250–290 °C, is one of the most commonly and widely used engineering plastic with various applications in automotive, electronics, construction and textile [6]. Therefore, research is being done on how to successfully incorporate natural fibers into polyamide without degrading it or resulting in composites with poor properties. One way has been to introduce additives to lower the processing temperature of polyamide enough to incorporate natural fibers [7]. This involved a two-step process where the polyamide was first melt compounded with the additive and then reinforced with the natural fibers. The results were composites with superior mechanical properties in comparison to the neat polyamide. However, the added processing step and use of expensive as well as non-environmental friendly additive increases the production cost of the composites, therefore making it economically and environmentally not feasible. Another method that was used to tackle this problem was to incorporate the filler at a later stage of the melt compounding process towards the end of the extruder [8]. By doing this, the fillers can be used without significant degradation occurring. The disadvantage to this process is that the short processing time doesn’t allow for proper dispersion of the filler in the matrix and therefore results to a composite with inconsistent dispersion of filler. Consequently, this leads to localized stress concentrations within the composite and premature failure. Solution blending of fibers with polyamide is another method that has been used to avoid the degradation of the fibers [9]. In this method, the polyamide was made into a solution while cellulose was dispersed in another liquid which is miscible with the polyamide solution. These two were then mixed together. Composite fibers were then made from the mixture using an injection spinning device. The merit to this process is that the filler is better dispersed within the polyamide with the added effect of no degradation. However, the process of fabrication involves the use of harsh and non-environmentally friendly chemicals, multi-step processes which incur costs and most importantly, is not industrially feasible.

As the result of the above tried methods, thermally treated torrefied biomass was added to polyamide for the fabrication of composite materials [10]. It was found that the torrefied biomass was more thermally stable than using the parent biomass directly as the filler. The absence of odor during melt compounding and improvement in mechanical properties in comparison to the untreated biomass composite suggested that its increased thermal stability led to better structural integrity of the composites. However, the presence of unevenly torrefied biomass led to some degradation during composite fabrication and ultimately reduced the mechanical properties with increasing filler loading. It is therefore of great importance to achieve an even thermally stable filler. Recently, biocarbon, a thermally treated biomass under limited or complete absence of oxygen has come under scrutiny as potential candidate for many applications. It is the by-product of the bio-refinery industry and has mainly found uses as solid fuel in the same industry and soil amendment in agriculture [11, 12, 13]. Its attraction due to renewability, sustainability, low cost and environmental friendliness through carbon sequestration has gained the attention of material scientists to be used as a possible alternative for value added purposes. Biocarbon is thermally stable enough to be used as reinforcing filler in plastics especially in engineering plastics without the addition of additives or limiting the exposure time to high processing temperatures. Biocarbon has a thermally stable structure because of its pyrolysis temperature which is typically above 350 °C [14]. It also has functional groups on its surfaces which can act as anchor points for polymer side and end groups [15]. Being relatively new filler to the field of polymer science and composites, it is of great importance, as a first step, to study the effects of biocarbon in various polymeric systems in fabricating composite materials.

As a result, an extensive effort has been put into the investigation of biocarbon as reinforcing fillers in many thermoplastics such as PLA [16, 17], rubber [18, 19, 20, 21], PP [22] and thermosets [23, 24, 25, 26]. These works studied the effect of biocarbon on the various properties of the plastics such as mechanical and thermal for example. They also investigated how the biocarbon concentration influenced these properties. In our previous studies, the preparation of biocarbon based polyamide bio-composites was also investigated [27, 28] but focused on the effect of particle size, pyrolysis conditions and durability of biocarbon on the properties of polyamide composites. The effect of biocarbon on the water uptake of polyamide was also investigated showing the benefit of using biocarbon [29]. However, no work till date has investigated the effect of varying biocarbon to a high concentration on the thermo-mechanical properties of polyamide. Typically, addition of natural fibers above 20–25 wt% results in increased volatization of the fiber components which lead to composites with poor properties. In this study, the effect of biocarbon on polyamide up to 40 wt% concentration to determine it viability as a bio-based reinforcing filler at high concentrations was investigated and reported. To achieve this, the mechanical, morphological and thermal properties were studied. For the mechanical properties, unidirectional tensile tests, three-point bend test and impact testing were performed. Differential scanning calorimetry, rheology and thermogravimetric analysis were done to determine the thermal properties. Scanning electron microscopy was used to characterize the morphology of the impact fractured surface. The results from this study act as a strong and vital background to future and further studies on biocarbon reinforced polyamide composites.

Materials and Methods


Polyamide 6 used in this work is heat stabilized Ultramid B3K in pellet form with a density of 1.13 g/cm3 and manufactured by BASF, USA. Miscanthus fibers were obtained from New Energy Farms Leamington, Ontario, Canada and pyrolyzed at 650 °C by Competitive Green Technologies, Leamington, ON, Canada, after which it was hammer milled to a particle size of ≤ 397 µm.

Sample Preparation and Compounding

The biocarbon was milled using procedures previously described [30]. Prior to processing, it was dried at 105 °C in a conventional oven until constant weight while the polyamide was dried at 80 °C overnight. Biocarbon reinforced polyamide composites were made using biocarbon loading at 5, 10, 20, 30 and 40 wt% in a micro compounder (DSM Xplore, Netherlands) equipped with a co-rotating twin screw and barrel volume of 15 cc. A processing time, temperature and screw speed of 120 s, 250 °C and 100 rpm were used for all composite fabrications respectively. The molten composites were collected and injection molded into test samples using a transfer device and a micro-injection molder (DSM Xplore) kept at 245 °C with the mold temperature set at 70 °C.

Morphological Analysis

The morphology of the impact fractured surfaces of the composites was examined using a scanning electron microscopy (SEM, FEI Inspect S50, OR, USA). All samples were sputter-coated twice with gold under vacuum for 60 s using a Cressington Sputter Coater 108 auto (Cressington Scientific Instruments, UK).

Thermal and Thermomechanical Analysis

The thermal properties such as the melting temperature (Tm), melt crystallization temperature (Tmc) and crystallinity (χc) were studied using differential scanning calorimeter (DSC Q200), manufactured by TA Instruments, USA, under nitrogen flow at a rate of 50 mL/min. All tests were done according to ASTM standard D7426. The melting enthalpy of fully crystalline polyamide used to calculate the degree of crystallinity was ΔH0m = 190 J/g. Injection molded samples ranging between 5 and 10 mg in weight were placed in an aluminum pan and heated from 23 to 250 °C at a rate of 10 °C/min, held at 250 °C for 5 min, cooled to 23 °C at a rate of 10 °C/min, held at 23 °C for 5 min and then heated again to 250 °C at the rate of 10 °C/min. The possible interactions occurring within the composites were investigated by studying the viscoelastic behavior of the samples using dynamic mechanical analyzer (DMA) (Q800, TA Instruments) in three-point bend clamp mode with a frequency of 1 Hz and a strain amplitude of 25 µm. Testing was done within a temperature range of − 10 to 150 °C at a ramp rate of 3 °C/min and a pre-load of 1 N. Samples with an average width and thickness of 12.6 and 3.35 mm were used respectively.

The thermal stability of polyamide and its composites was measured by means of a Thermo Gravimetric Analyzer (TGA) (Q500, TA Instruments). For each test, 15–20 mg of material was used. All tests were done within a temperature range of 23–600 °C, ramp rate of 20 °C/min and in either a nitrogen or air environment with a flow of 40 mL/min. For the calculation of the activation energy of thermal degradation kinetics, the experiments were also conducted at 5 and 10 °C/min.

HDT measurement was performed based on the ASTM D648 standard at a constant load 0.455 MPa in the same DMA instrument. The analysis was performed at a heating rate of 2 °C/min−1 from ambient temperature to 100 °C in a three-point bending mode. Two samples have been tested for each formulation.

Mechanical Properties

All samples were conditioned at 23 ± 2 °C and 50 ± 5% relative humidity for a period of 40 h before testing. Tensile properties were measured using samples prepared according to Type IV specimen of ASTM standard D638. The tests were conducted on a Universal Testing Machine (Instron, Norwood, MA, USA) under room temperature at a speed of 50 mm/min for both neat polyamide and its composites for purpose of direct comparison to the neat polyamide. However, this speed used for testing the composites does not confirm to ASTM standard. Flexural properties were measured using samples prepared according to procedure B of ASTM standard D790. The tests were conducted on a Universal Testing Machine (Instron, Norwood, MA, USA) under room temperature at a crosshead speed of 14 mm/min and 23 °C. The impact strength of the samples was measured using samples prepared in accordance with ASTM standard D256 with notched depths of 2 mm made using a TMI Notching Cutter. The tests were conducted on a TMI Monitor Impact Tester (Testing Machines Inc., DE, USA) at room temperature.

Melt Flow Index Analysis

Melt flow index of the neat polyamide and its composites was measured according to the ASTM D1238 standard using a Melt Flow Indexer (Qualitest model 2000A) at 235 °C with a standard weight of 2.16 kg. For each measurement, approximately 6 g of the material was used and the first cut during measurement taken. The measurement of each sample was repeated five times to obtain five replicates. The results are an average of the five replicates for each sample.

Rheological Analysis

Anton Paar Modular Compact Rheometer MCR-302 was used to measure the shear and complex viscosities, storage modulus and tan delta of polyamide and its composites. Injection molded disks samples were placed between 25 mm parallel plates and measured at 240 °C and at a thickness of 1 mm.

Results and Discussion

Differential Scanning Calorimetric Analysis

The thermal properties from the DSC analysis are given in Table 1. The results show that the melt crystallization temperature (Tmc) of polyamide decreased slightly with the addition of biocarbon. At 40 wt% biocarbon loading, the Tmc was reduced by approximately 3 °C. This suggests that the biocarbon plays a role in preventing the nucleation of crystals in the polyamide, therefore shifting the temperature to a lower value. By integration of the melting peak, the degree of crystallinity was calculated taking into account the weight fraction of polyamide in the specimens processed. The degree of crystallinity of the polyamide was observed to reduce by approximately 68.6% when 40 wt% biocarbon was used. The reduction in the crystallinity of the composites is expected to have a significant effect on the moduli as the moduli of polymers are related and dependent on the degree of crystallinity.

Table 1

Data from the DSC analysis for polyamide and biocarbon based bio-composites

Polyamide/biocarbon (wt/wt) %

Tmc (°C)

ΔHm (J/g)

Tm (°C)

χc (%)


192.8 (0.2)

62.95 (1.15)

220.5 (0.2)

33.13 (0.61)


192.3 (0.2)

57.04 (1.00)

220.4 (0.1)

28.52 (0.50)


191.6 (0.1)

54.18 (0.94)

220.6 (0.4)

25.66 (0.45)


191.4 (0.2)

44.20 (1.60)

221.0 (0.4)

18.61 (0.67)


190.1 (0.1)

38.79 (1.89)

220.2 (0.1)

14.29 (0.70)


189.6 (0.1)

32.98 (0.72)

220.4 (0.2)

10.41 (0.23)

Thermogravimetric Analysis

Thermal Stability Analysis

The thermal stability of polyamide and its composites was analyzed by TGA under either nitrogen (Fig. 1A) or air (Fig. 1B) atmospheres. The figures show the percentage weight loss (solid lines) and the derivative thermogravimetric thermograms, DTG, (dash lines) as a function of temperature. The specimens were characterized by their temperature at 10% weight loss (T10%) and temperature of maximum degradation rate, determined by the position of the maximum on the differential TGA curve (Tmax); all the temperatures are summarized in Table 2.

Fig. 1

Thermogravimetric analysis TGA and DTG in (A) nitrogen and (B) air environment of polyamide 6 and its composites at different biocarbon loading

Table 2

TGA results for polyamide and biocarbon based bio-composites in either nitrogen or air flow

Polyamide/biocarbon (wt/wt) %

TGA under air flow

TGA under nitrogen flow

T10% (°C)

Tmax (°C)

T10% (°C)

Tmax (°C)































The direct comparison of TGA results obtained in nitrogen and oxygen atmospheres highlights that the specimens undergo a two steps decomposition. In air atmosphere, only the first weight loss is observed on both the polymer matrix and bio-composites. In the range of temperatures investigated, only the broad plateau of the second degradation step is observed and the consecutive main weight loss is expected to occurs at higher temperature (between 600 and 800 °C). In nitrogen flow, the second step is observed on the thermogram of polyamide as a simple shoulder while the bio-composites show a significant degradation peak which can be attributed to the carbonization processes of the materials. This second degradation process can be observed to increase with the increase of biocarbon content. The nature of the flow applied further plays a role in the polyamide degradation. In air, the neat polyamide degradation starts (T10%) at 412 °C; a shift to higher temperature (417 °C) is observed when using nitrogen (Table 2). The lower value observed in air is attributed to the association of thermal and oxidative degradations occurring in presence of oxygen while thermal degradation only occurs in inert nitrogen. This reactive environment induces more differences than in a nitrogen environment. In air, the onset of degradation (T10%) of polyamide shifts from 412 to 424 °C (a 12 °C increase) at 40 wt% loading, with a main shift occurring for 20 wt% biocarbon. Likewise, in nitrogen it shifts from 417 to 426 °C (a 9 °C increase) with a big shift after 40 wt% loading (Table 2). Regarding the Tmax, a shift to lower temperature from 467 to 458 °C (a 9 °C reduction) is observed in air and significantly decreasing for 40 wt% biocarbon loading. Whereas, in nitrogen, the decrease observed occurs gradually from 466 to 459 °C (a 7 °C reduction). Thus, the biocarbon addition delays the degradation onset of polyamide but increases the degradation rate. The improved thermal stability observed after biocarbon addition can be attributed to a shielding effect of biocarbon. This shielding or barrier effect can take different forms. One way is chemically by hindering volatilization and diffusion of decomposed volatiles products. The other way is physically by improving the thermal hysteresis of the composites i.e. decreasing the heat transfer of the material by adding filler with a higher heat capacity or by restricting the molecular motion of the polymer [31]. It can be observed that lower amount of filler i.e. below 20 wt% has no influence on either T10% and Tmax in both atmospheres investigated. This can be explained by the heterogeneous effect of biocarbon micro-particles added in too low amounts. Sufficient particle content is required to get a homogenous dispersion which can then effectively act as a shield to the underlying polymer matrix. The heat capacity measurement of both PA6 and biocarbon in the melt could partially explain the difference obtained. At 25 °C, the matrix shows a heat capacity (Cp) value of 1449 J/kg/K [32] while biocarbon from miscanthus has been shown to have a Cp value of 2300 J/g/K [14]. It is expected that a rise in temperature would increase these values, but if such difference is maintained at the degradation temperature then the biocarbon particles would act as a shield by absorbing the heat which would preserve polyamide chain degradation; this behavior accented by increasing the heating rate. However, considering the difference in trends observed (brutal shift at 20 and 40 wt% biocarbon content in air and nitrogen respectively), if either of the physical effects was responsible for the shielding effect, then the same trend would be observed in whatever the atmosphere is. The difference observed between T10% in air and nitrogen points to the dominant effect of the volatiles absorption by the biocarbon particles, increasing significantly the thermal stability. The presence of functional groups on the biocarbon surface [27] and the high specific surface area induced by the particles porosity [14] would absorb the volatile polar products [33] given off during the degradation of polyamide. In air, the shielding effect occurring above 20 wt% biocarbon content shows a shift of the degradation onset from approximately 410 °C to a range between 420 and 425 °C, stressing a similar effect as using a nitrogen atmosphere (Table 2). Moreover, the significant shift to higher temperature observed for T10% in nitrogen flow at 40 wt% loading would highlight the influence of another effect, which is probably due to the percolation of biocarbon particles (this effect is developed in more details later). The shift to lower temperature obtained for Tmax was further developed in the next thermal degradation kinetic section.

The TGA curves also help to estimate the real biocarbon content of the bio-composite. At 500 °C, the polyamide has completely degraded leaving nothing but 0.39% residual weight. Therefore, the residual weights of all the bio-composites were determined at 500 °C. It was determined that the 5, 10, 20, 30 and 40 wt% biocarbon filled composites had 5.95, 8.90, 19.11, 29.17 and 39.08 wt% residual weights respectively. This result shows the precision of the biocarbon content within the polyamide matrix.

When we observe the derivative thermogravimetric (DTG) curves of the samples in nitrogen in Fig. 1A, we can see that although the thermograms are slightly shifted to lower temperature, the shape is similar though proportional to the filler loading. This behavior suggests that the rates of degradation are comparable between all the samples. In opposition, the samples tested in air environment show a slight reduction in degradation rate after biocarbon addition. To better understand the effect of biocarbon and its concentration on the degradation behavior of polyamide, the activation energies of all the samples were determined and analyzed.

Activation Energy of Thermal Degradation Kinetics

The effect of biocarbon loading on the degradation of polyamide was further analyzed by calculating the activation energy of the samples at different heating rates i.e. 5, 10 and 20 °C/min in both nitrogen and air. The general form of the basic rate equation is given by the Arrhenius law (Eq. 1) representing the relationship between process rates and parameters such as temperature (T) and extent of conversion (α):
$$\frac{{{{\text{d}{\upalpha}}}}}{{{\text{f}}({{{\upalpha}}})}}=~\frac{{\text{A}}}{{\text{q}}}{\text{exp}}\left( {\frac{{ - {\text{E}}}}{{{\text{RT}}}}} \right){\text{dT}}$$
where α and f(α) are the conversion degree and reaction model respectively. q, E, R and A are the heating rate, activation energy, universal gas constant and pre-exponential factor respectively. The conversion degree of the thermal degradation at time t is obtained from Eq. 2.
$${{{\upalpha}}}=\left( {\frac{{{{\text{m}}_{\text{i}}} - {{\text{m}}_{\text{T}}}}}{{{{\text{m}}_{\text{i}}} - {{\text{m}}_{\text{f}}}}}} \right)$$
where mi and mf are the initial and final masses while mT is the mass at a specific temperature T. Equation 1 can further be approximated using Doyle’s approximation to derive Eq. 3 which plots the log(q) versus inverse of temperature for a given α. This produces a straight fitted line which can then be used to estimate the activation energy at different α from the regression of the line.
$$~{\text{log}}({\text{q}})={\text{ constant }}-0.{\text{4567}}\left( {\frac{{\text{E}}}{{{\text{RT}}}}} \right)$$

From Eq. 2, the plots of log(q) versus inverse of temperature for different given α of the samples is given in Fig. 2A, B. It can be observed that the iso-conversion of all the samples at different α fit well to a straight line especially for those performed under nitrogen environment.

Fig. 2

Iso-conversional plots of biocarbon filled polyamide 6 composites performed under (A) nitrogen and (B) air environments

Figure 3A shows the plot of the activation energy (E) versus α in nitrogen atmosphere. It can be observed that the activation energy clearly decreases with the addition of biocarbon to polyamide. The activation energy for the polyamide can be observed to be stable and slightly increases with increasing conversions within a range from 180.9 to 191.2 KJ/mol. Similar values were observed for polyamide 6 to be at 180 ± 10 and 201 KJ/mol in nitrogen atmosphere, showing the high consistency of the data [34, 35]. With the addition of biocarbon (5 wt%), a significant decrease in E was observed which further decreased with increasing conversion. At high biocarbon loadings, we observe some increase in the activation energy, however, always lower than that of the neat polyamide. This result would indicate that the increase in thermal stability observed after biocarbon addition cannot be attributed to a higher energetic barrier. Accordingly, biocarbon would act also as a catalyst to aid molecular break down of the polyamide chains. While the presence of elements (alkali, alkaline earth and transition metal) is known to catalyze the biocarbon conversion [36] and could affect the polyamide degradation, none of these elements was used for the biocarbon preparation and no evidence of their presence was highlighted from energy-dispersive X-ray spectroscopy (EDS). However, during the thermal decomposition of the biocarbon, volatiles are released which can act to catalyze the decomposition of the polyamide chains. Biocarbon has been investigated previously as a potential catalyst in the decomposition of toluene and tar [37, 38]. Therefore, it is not surprising that biocarbon can also catalyze the decomposition of polymer chains. Amintowlieh et al. [35] found that addition of wheat straw to polyamide 6 significantly reduced the thermal stability and activation energy of the bio-composites. However, biocarbon is more thermally stable with a lot less thermally unstable components than in natural fibers which decompose at low temperatures. Hence, resulting to better thermal stability of the bio-composites when compared to natural fiber reinforced polyamide.

Fig. 3

Activation energy of biocarbon reinforced polyamide 6 at different loadings and at different conversions in (A) nitrogen and (B) air environments

From the activation energy of the samples in air as seen in Fig. 3B, air has a significant effect on the degradation kinetics of the samples. The activation energy shows a completely different trend in comparison to that of samples tested in a nitrogen environment; especially for the bio-composites. The E are increased significantly with the bio-composites having higher values across the different conversions except for 5 wt% biocarbon reinforced polyamide when compared to that of the neat polyamide. This result for the most part shows that in the presence of an oxidizing environment, biocarbon can increase the activation energy of polyamide. This can further be observed from the 10 wt% biocarbon reinforced polyamide sample which is the most significantly improved of all the samples. This result is consistent with the theory proposed earlier. Likewise, at conversions higher than 0.5, samples 10, 20, 30 and 40 wt% biocarbon reinforced polyamide exhibit maximum activation energies indicated with red arrows in Fig. 3B and drastically reduce thereafter. These maximum energy peaks are attributed to the biocarbon carbonization as they occur for high degree of conversion and specifically at the biocarbon content. This would indicate that biocarbon is carbonizing at a higher energetic barrier and once the degradation is initiated, the energy decreases. It is impossible to ascertain what occurs during the degradation of polyamide in the presence of biocarbon and air. This is because of the multiple reactions occurring and the presence of multiple gases which can influence the reactions. However, from the addition of biocarbon to polyamide in the presence of an oxidizing environment, greater energy is required to degrade the molecular interactions occurring in the samples indicating the possible formation of stronger and more stable bonds with the biocarbon or crosslinking of the polyamide chains. Similar results have been found for polyamide 6/clay nanocomposites [39]. The activation energy of the polyamide was improved in the presence of clay and air.

Mechanical Properties

Tensile and Flexural Behaviors

The tensile and flexural strengths of polymer composites are highly dependent on the effectiveness of the adhesion between matrix and filler which leads to efficient stress transfer from matrix to filler. As shown elsewhere, a high interaction between polyamide and biocarbon was observed by morphological investigation though not evidenced by the detection of any strong chemicals bonds. The high affinity between polyamide and biocarbon was attributed to van der Walls and hydrogen bonding between the amide group of polyamide and the functional groups of biocarbon [28]. It can be observed from Fig. 4 that the tensile and flexural strengths follow similar trends; they increase with increasing biocarbon content. They also show different transitions regions; 1, 2 and 3 as labelled on the graph. The transition regions are more significant and distinct with the tensile strength than with the flexural strength. The flexural strength is significantly higher than the tensile strength because of the added region of resistance to compression during bending tests. The biocarbons restrict the polyamide chains from deformation. In region 1, we observe an insignificant increase in the strength with the addition of 5 wt% biocarbon. Likewise, in region 2, there is little to no significant increase as well despite the addition of 10 and 20 wt% biocarbon. However, in region 3 we observe a sudden and significant increase in strength. The change in slope in region 3 after 20 wt% biocarbon addition is most likely due to a sudden increase and more efficient stress transfer. One reason for efficient stress transfer could be as a result of the reduction in the inter-particle distance. As particles come closer to each other, stress transfer efficiency is increased as stress can easily and rapidly be transferred from particle to particle. This is because of the percolation effect where particles come close to or begin to come in contact with one another. At 40 wt% biocarbon loading, the tensile and flexural strengths are enhanced over the neat polyamide by 20 and 47% respectively.

Fig. 4

Effect of biocarbon loading on the tensile and flexural strengths of polyamide 6

To understand the behavior of the tensile strength, the tensile fractured surfaces of the bio-composites were analyzed employing SEM and the images are shown in Fig. 5A–F. From Fig. 5, the neat polyamide shows a flat and brittle mode of fracture with ridge-like grooves after work hardening of the sample. However, upon adding 5 wt% of biocarbon, the fracture mode is changed to the elongation of the matrix ligaments and fibril fracture. It can be seen from the zoomed-in morphology that some of the biocarbon particles are surrounded by the fibrils while still connected by polyamide ligaments as indicated by circles. Similar structure is observed in the morphology of 10 wt% biocarbon filled polyamide. This indicates that the biocarbon and polyamide have good affinity and therefore bound to one another. However, there are also particles which show voids around the interfaces with the polyamide. This difference in wetting might have originated from the pyrolysis and milling process of the biocarbon prior to composite fabrication. When the fibers are pyrolyzed, the surfaces exposed to the heat and limited oxygen environment oxidize to form surface functionalities while the core remains inert. During milling, the pyrolyzed fibers are crushed, exposing new and un-oxidized surfaces with no surface functionalities. This results in particles with both functionalized and un-functionalized surfaces, therefore leading to good and poor adhesion with the polyamide. Upon observing the morphology of the 20 wt% filled polyamide, the fracture mode changes with reduction in fibrillation of the polyamide ligaments while also showing more particles well bonded to the polyamide. At 30 wt% biocarbon filled polyamide, similar fracture mode is observed but with smaller ligament thickness due to the presence of more particles. We also observe even more particles that are well adhered to and embedded in the polyamide. The morphology of the 40 wt% filled polyamide shows a completely different fracture mode with no fibrillation occurring; a rough surface with well wetted and embedded biocarbon particles is observed.

Fig. 5

Micrographs of the tensile fractured surfaces of (A) 0 wt%, (B) 5 wt%, (C) 10 wt%, (D) 20 wt%, (E) 30 wt% and (F) 40 wt% biocarbon content in polyamide 6

The morphological analysis from Fig. 5 revealed that below 20 wt% biocarbon loading, stress is mostly carried by the matrix because of the presence of few well bonded biocarbon particles with surface functionalities. Beyond that, we observe significantly more bonded particles to the polyamide due to increase in particles with more surface functionalities. This leads to increased stress bearing by the biocarbon particles, therefore translating to significant increase in strength.

Like the tensile and flexural strengths, the moduli also follow similar trends; an increase with the addition of biocarbon and a positive transition in slope at 20 wt% (Fig. 6). It is very typical for a filler to increase the stiffness of a polymer with increasing filler content. This behavior is due to the significantly higher stiffness of biocarbon characterized by a Young modulus measured around 4–5 GPa after pyrolysis at ~ 500 °C, in comparison to polyamide 6 of 2.5 GPa [27]. This happens when the filler restricts the chain mobility of the matrix and therefore prevents it from easily deforming. The increase observed at and after 20 wt% could be related to the percolation of particles which cause further restrictions and therefore increase the stiffness of the composite. Another interesting aspect to the tensile moduli of the composites is that the increase observed is not as drastic as those of the flexural strength. This can be attributed to the decrease in the crystallinity of the composites. Though the biocarbons are stiffer than the polyamide and restrict the mobility of its chains, the reduction in the crystallinity of the polyamide affects its stiffness and therefore has a counteracting effect of the biocarbons. This leads to a gradual increase in the moduli. On the other hand, the flexural strength shows more significant increase because of the resistance to deformation of the polyamide chains in two ways; tension and compression when the samples are subjected to the bend test. The significant improvement of the mechanical properties after the addition of biocarbon can be attractive to many areas where high modulus and strength are required, like in automotive industry [40].

Fig. 6

Effect of biocarbon loading on the tensile and flexural strengths of polyamide 6

Impact and Elongation Behaviors

Impact strength and elongation typically depend on the plasticity of the polymeric material; its ability to plastically deform and absorb energy when some sort of stress is applied to it. Where there is an increase in elongation, the impact strength generally increases as well due to the ability of the matrix to plastically deform. However, in these bio-composites, the contrary is observed. The elongation at break is significantly hindered by the addition of 5 wt% biocarbon with a decrease of 74% (Fig. 7). Further addition of biocarbon up to 40 wt% showed further decreasing trend with a decrease of 88%. It was expected, as the biocarbon restricts the polyamide chains from moving or sliding during deformation. Likewise, the restricting effect of the polyamide chains was enhanced by the affinity between the polyamide chains and biocarbon particles [27]. However, the impact strength showed an increase with increasing biocarbon content with a maximum at 20 wt% biocarbon content of 33% increase. Beyond that we observed a decrease, and at 40 wt% biocarbon content, the strength was comparable to that of the neat polyamide.

Fig. 7

Effect of biocarbon loading on the impact strength and elongation at break of polyamide 6

Though the biocarbons restrict the polyamide chains from mobility, other mechanisms of impact toughening take place and inhibit the crack propagation. It has been shown that the addition of fillers to polymers can enhance the impact strength of a matrix up to a certain filler concentrations [41] through mechanisms such as crack pinning [28] and bowing by stopping or diverting the crack [42]. The fracture transitions from ductile to a brittle behavior in relation to filler concentration has been described in literature as the optimum toughening filler value above which there is no toughening observed [43]. This is probably as a result of percolation effect taking over with a transition from polymer-filler to filler–filler interactions leading to poor crack deflection and sudden failure [41].

Figure 8 shows the fractured surfaces of polyamide and its composites with biocarbon at different loadings. No clear difference could be observed from the surface topology except that biocarbon particles were exposed starting at samples containing 30 wt% biocarbon loading. This could be due to the crack propagating through particles which are percolating, thereby exposing them. It is expected that the impact strength of the bio-composite will be diminished since these points act as weak and easy points at which cracks can propagate through. To verify the percolation of the biocarbon particles at higher loadings, the bio-composites were microtomed and a cross-section of each of the samples was obtained and are shown in Fig. 9.

Fig. 8

Micrographs of the impact fractured surfaces of (A) neat polyamide and polyamide with biocarbon loadings of (B) 5 wt%, (C) 10 wt%, (D) 20 wt%, (E) 30 wt% and (F) 40 wt%

Fig. 9

Micrographs of the microtomed surfaces of (A) neat polyamide, (B) 5 wt%, (C) 10 wt%, (D) 20 wt%, (E) 30 wt% and (F) 40 wt% biocarbon content in polyamide

As observed from Fig. 9, with increasing biocarbon concentration, the inter-particle distance is reduced. Above 20 wt% biocarbon loading, the particle–particle contact or percolation effect could be observed. At this stage, the stress imparted to the surrounding polyamide is high and causes significant restriction to deformation and therefore resulting to poor impact strength. Equally and inversely, the same effect results to improvement in the tensile and flexural strengths as describe previously.

Heat Deflection Temperature (HDT)

The results from the HDT measurements of the samples are given in Table 3. Upon the addition of 5 wt% biocarbon to the polyamide, the HDT is increased. With addition of higher amounts of biocarbon to the polyamide up to 40 wt%, the subsequent increases in the HDT value was observed. This was due to the effect of the biocarbon on the polyamide which acted to restrict chain mobility, as shown elsewhere [27]. With a constant load being applied and rise in temperature, the presence of biocarbon in the polyamide prevented it from deforming under the load as the chains of the polyamide became more and more mobile, therefore, leading to a higher deflection temperature.

Table 3

Results of HDT, MFI and zero shear viscosity measurements of polyamide and its composites with biocarbons at different loadings

Polyamide/biocarbon (wt/wt) %

HDT (°C)

MFI (g/10 min)

Zero shear viscosity (Pa s)

R-squared (R2)


164.6 (3.1)

30.0 (0.3)




168.8 (0.2)

20.4 (1.2)




172.1 (1.0)

17.1 (0.6)




184.2 (3.3)

10.1 (0.1)




194.9 (0.1)

5.8 (0.5)




196.9 (4.4)

3.6 (0.0)



Melt Flow Index (MFI)

The MFI of the neat polyamide and its composites at various biocarbon loadings were measured and the results are given in Table 3. The addition of biocarbon into polyamide reduces the MFI: a 33 and 88% reduction for 5 and 40 wt% of biocarbon loadings were observed respectively. This is because of the inclusion of the biocarbons which restrict the flow of the polyamide chains, therefore increasing the viscosity of the samples. It is important for the composite to flow in melt state to easily fill molds during part production. An MFI value of about 10 g/10 min is typically good enough to fill molds with thick walls. It should be noted that these measurements were done at 235 °C with a load of 2.16 kg, to satisfy the standard of polyamide. Therefore, during extrusion and injection molding, the flow index will be much higher than the values obtain here as injection molding of polyamide and polyamide composites are typically done at higher pressures and temperatures.

Rheological Measurements

Like the MFI, the viscosity is a measure of the polymers response by flow under applied stress or force and provides important information about processing and performance. The average of three viscosity tests per sample of the neat polyamide and its composites were measured and reported in Table 3 and Fig. 10. From Fig. 10A, it can be observed that there are 2 regions of somewhat stable viscosities in the polyamide. These regions are in the ranges of 0.001–0.1 and 0.1–10 s−1. This indicates that at higher shear rates (around 0.1–1 s−1), there is some disentanglement of the polyamide chains leading to increased flow. This behavior can also be noticed for the composite samples as well. However, there is a slight but progressive shift of the transition between the regions to lower shear rates with increased biocarbon loading. This implies that the biocarbon is obstructing chain entanglement and therefore shifting the shear thinning to lower shear rates. Also, across the span of shear rates, the viscosity is increased with increasing biocarbon loading. Similarly, as observed earlier, the increase in moduli and reduction in crystallinity show the strong effect that biocarbons have on the polyamide. Tim et al [44] also found an increase of viscosity when biocarbon was added to polyamide.

Fig. 10

(A) Viscosity versus shear rate and (B) Carreau-Yasuda model of neat polyamide and its composites

The viscosity of any polymer is always measured at specific shear rates. However, to get the viscosity of a polymer when the chains are at rest and not slipping past each other due to the applied shear, the Carreau-Yasuda model can be used to extrapolate the zero shear viscosity (ZSV) by modelling the experimental viscosity data of polymers. The ZSV can be calculated from the Carreau-Yasuda model using Eq. (4):
$$\eta \left( {\dot {\gamma }} \right)={\eta _\infty }+({\eta _0} - ~{\eta _\infty }){(1+{(\lambda \dot {\gamma })^a})^{\frac{{n - 1}}{a}}}$$
where \(\eta\), \({\eta }_{0}\), \({\eta }_{\infty }\) and \(\dot{\gamma }\) correspond to the viscosity values, ZSV, viscosity at infinity and shear rate respectively. a, n and λ are empirically determined constant parameters having no units except for λ which is in seconds. The curves in Fig. 10A were used to generate the Carreau-Yasuda model through the Rheoplus software and are represented in Fig. 10B. The ZSV for the samples are given in Table 3 as generated by the software. The ZSVs are higher with the addition of biocarbon. At 30 wt% loading, the ZSV is greatly increased, by 2.6 times, in comparison to that of 20 wt%. It could suggest that at 30 wt%, the polyamide chains motions are highly restricted because of the increase in biocarbon particles, therefore having closer inter-particle distances. The R-squared (Table 3) tells how well the model fits the experimental data. As it can be seen, a high percentage indicates that the model fits the data well. The 20 and 40 wt% filled polyamide composite samples showed a lower R-squared value than other samples. Regardless, it is still high and therefore indicative of a strong fit. However, it could be observed for these concentrations, there are slight differences in behaviors due to the specific concentrations of biocarbons. It has been shown in the previous properties that 20 wt% filled polyamide had the highest impact strength and was the concentration at which a transition to a steeper moduli and strength occurred.


Biocarbon obtained using miscanthus fibers was effectively used as a reinforcement in polyamide at different loadings up to 40 wt%. It was observed that the biocarbon had very strong interactions with the polyamide chains resulting in great chain restrictions. Subsequently, the moduli were observed to increase as a result. The HDTs of the composites were also observed to increase with increasing biocarbon content. Biocarbon unlike natural fibers do not diminish the thermal stability of the resulting bio-composites as much. However, in a nitrogen environment, the presence of volatiles stemming from the surface functionalities still cause some reduction in the thermal stability of the polyamide 6. However, in an oxidizing environment such as air, the biocarbon is observed to promote thermal stability by increasing the activation energy of the polyamide. Therefore, it can be said that depending on the environment, biocarbons can either act as an anti-oxidizing agent of catalyst for degradation. The tensile and flexural strengths were greatly improved because of good adhesion between the biocarbon and polyamide matrix which was observed from the morphology of the impact fractured surfaces of the composites. The impact strengths of all the composites were comparable to that of the neat polyamide despite the incorporation of biocarbon. However, at 20 wt% biocarbon loading, it was observed to increase above that of the neat polyamide. The MFI and viscosity of the polyamide was decreased and increased respectively with the addition of biocarbon. The rheological analysis suggests that the presence of biocarbon in the polyamide increases chain disentanglement. This behavior is increased with increasing biocarbon content as well. The fabrication of biocarbon reinforced polyamide composites show promise to be used in automotive underhood applications such as air intake manifolds and electric circuit casing, where mineral filled polyamide composites are applied as well as composites which require high filler loading to offset the cost of the matrix while still maintaining some level of flexibility.



The authors acknowledge the financial support by the Ontario Ministry of Agriculture, Food and Rural Affairs (OMAFRA) – University of Guelph Product Development and Enhancement through Value Chains Research Theme (Project # 200399, 200388, 200245), The Natural Sciences and Engineering Research Council (NSERC), Canada Discovery grant (Project # 400322) and Ontario Research Fund, Research Excellence Program; Round-7 (ORF-RE07) from the Ontario Ministry of Research and Innovation, currently known as the Ontario Ministry of Research, Innovation and Science (MRIS) (Project # 052644 and # 052665).


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Copyright information

© Springer Science+Business Media, LLC, part of Springer Nature 2018

Authors and Affiliations

  • Emmanuel O. Ogunsona
    • 1
    • 2
  • Amandine Codou
    • 2
  • Manjusri Misra
    • 1
    • 2
  • Amar K. Mohanty
    • 1
    • 2
  1. 1.School of EngineeringUniversity of GuelphGuelphCanada
  2. 2.Bioproducts Discovery and Development Center, Department of Plant AgricultureUniversity of GuelphGuelphCanada

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