Introduction

The growing global problem of plastic pollution has generated much interest in biodegradable green polymers for many industrial applications [1]. Polylactic acid (PLA) is a semicrystalline polymer which is further biodegradable and can be obtained from renewable agricultural products such as crop straw [2, 3]. Another advantage of PLA is that its decomposition products have no obvious harm to the environment. In addition, PLA also has excellent tensile strength, highly biocompatible, and high transparency. Therefore, PLA is regarded as a potential replacement of petro-derived polymers. Regrettably, PLA also has its own limitations in terms of poor thermal stability, low crystalline, and low crystallization rate, which greatly restrict its applications [4]. In particular, its low crystallization rate is still one of the most pressing technical problems of PLA industry applications. The crystallization behavior of PLA has an important effect on its mechanical, thermal stability as well as rheological behavior properties. Therefore, improving the crystallization behavior of PLA is a key factor to enhance the properties of PLA. The addition of nucleating agent can provide abundant nuclei in PLA matrix, which can affect the crystallization rate and crystallinity of PLA significantly. So, for a improve crystallinity of PLA the employment of nucleating agent can play an important role. At present, nucleating agents used to improve PLA crystallization mainly include organic nucleating agents and inorganic nucleating agents [5, 6]. Organic nucleating agents have received considerable attention due to their excellent nucleation. Jing et al. [5] studied the effect of N-Aminophthalimide compound (NA-S) as nucleating agent on crystallization behavior and morphology of PLA, in which the addition of 1% NA-S significantly improved the crystallization rate and crystallinity of PLA. Tang et al. [7] investigated the effect of ethylenebishydroxystearamide (EBH) as a crystal nucleating agent on the crystallization behavior of PLA. Results showed that the crystallization rate and mechanical performance of PLA were improved. Suksut et al. [8] reported that natural rubber improved the physical properties of PLA. Compared with organic nucleating agents, inorganic nucleating agents have insufficient nucleation effect, which can not be accepted by industrial production. Moreover, inorganic nucleating agents are generally incompatible with organic non-polar polymer, which is not conducive to improve the performance of PLA materials. Therefore, it is an urgent problem to find new and efficient nucleating agents. Currently, the most promising organic nucleating agents are the amide compounds. Song et al. [9] studied the effect of multiamide nucleator (TMC) on crystallization behavior of PLA, showed an evident increase in crystallinity and crystallization rate. Crystalline performance of PLA/1,3,5-benzenetricarboxylamide composites, investigated by Nakajima et al. [10], demonstrated an obvious increase in crystallization enthalpy and crystallization rate. Xing et al. [11] investigated the affects of aliphatic amides nucleating agent on crystallinity of PLA, which showed an enhanced crystallinity behavior of PLA. Cai et al. [12] studied the effect of bisamide nucleating agents on the thermal behavior of PLA. Results showed that the thermal stability of PLA was improved. Hence, a nucleating agent prepared by amidation, can be used to increase the crystallization rate of PLA and can thus improve the properties of PLA.

Layered double hydroxides (LDHs) are one kind of special class nano-sized lamellar compounds, which consist of a stacking of positively charged hydroxide layers with intercalated anionic species and water molecules [13]. The general formula of LDHs can be written as [MII1-x MIII x (OH)2x] x+(An-)x/n⋅ mH2O, where MII and MIII represent divalent and trivalent metal ions, respectively. and An- represents an anion. Due to their excellent anion exchange properties and highly adjustable properties, LDHs have been widely used as adsorbents, catalysts, flame retardants [14,15,16]. Moreover, there are many other advantages of LDHs, such as a simple and economical synthetic method, and are also nontoxic and biocompatible to human body. Hence, as a class of emerging layered materials, LDHs are considered to be of great value for the preparation of polymeric nanocomposites [17, 18]. However, the electrostatic interaction between the hydroxide interlayers and intercalated anions/water molecules of LDHs is so strong that the exfoliation or intercalation in polymer matrix is very difficult [19]. In addition, the interfacial compatibility between neat inorganic LDHs and polymer matrix is poor, which leads to a decline in its nucleation effect. Therefore, it is necessary to carry out organic modification of neat inorganic LDHs in the preparation of polymer/LDHs nanocomposites. So far, there are many studies focused on organically modified LDHs via intercalation method, aimed at achieving a good dispersion and exfoliation of LDHs in polymer matrix. For example, Costache et al. modified LDHs intercalated with sodium 4-styrenesulfonate (StSA) to give MStSA, which showed good dispersion in poly (methyl methacrylate) [20]. Recently, in order to widen the range applications of LDHs, there has been a growing interest in their use as nucleating agent in PLA materials to improve their physical properties [21]. Many studies involving PLA/LDHs nanocomposites and characterization of their different properties such as, mechanical, morphological, thermal, degradation stability etc. For examples, Teh et al. [22] investigated the effect of layered double hydroxides on the mechanical and morphological properties of poly (hydroxybutyrate)/poly (lactic acid) (PLA) nanocomposites, which showed a lower molding cycle time and better mechanical and morphological properties. Chiang et al. [23] studied the degradation stability of poly (lactic acid) (PLA)/γ-polyglutamate-modified layered double hydroxide nanocomposites, which showed an enhanced thermal degradation behavior of PLA nanocomposites.

The present study focuses on a novel way of developing new and highly efficient nucleating agents. To our best knowledge, there is no report on the use of synthetic amidated LDHs as a nucleating agent to improve the performance of PLA. In this paper, the organic intercalation modification of LDHs (E-LDHs) was carried out by EDTA. Then, the organic intercalated LDHs (E-LDHs) was reacted with aniline to obtain the amide EDTA intercalation LDHs (AE-LDHs). To improve the performance of PLA, PLA/AE-LDHs nanocomposites were prepared by melt blending. The structures and properties of PLA/AE-LDHs nanocomposites were also studied.

Experimental

Materials

PLA (injection grade, Mw = 100,000 g mol−1) was purchased from Shenzhen Guanghuaweiye Industrial Co. Ltd., Shenzhen, China. Zn(NO3)2⋅6H2O and Al(NO3)3⋅9H2O were all purchased from Fuchen Chemical Reagent Factory, Tianjin, China. Sodium hydroxide (NaOH) and sodium carbonate (Na2CO3) were obtained from Kelong Chemical Reagent Factory, Chengdu, China. N, N-Dimethylformamide (DMF, >99%), aniline (>99%) and toluene (>99.5%) were procured from Zhiyuan Chemical Reagent Co. Ltd., Tianjin, China. Ethanol solution (CH3CH2OH), boric acid (H3BO3), and ethylenediamine tetraacetic acid disodium salt (EDTA) were all purchased from Fuchen Chemical Reagent Factory, Tianjin, China.

Preparation of LDHs

LDHs was prepared by coprecipitation method. Zn(NO3)2⋅6H2O (17.82 g) and Al(NO3)3⋅9H2O (11.25 g) were dissolved in 100 mL of deionized water according to the molar ratio of n(Zn2+): n(Al3+) = 2:1. After stirring for some time, 2 mol/L NaOH solution and 1 mol/L Na2CO3 solution were slowly added into the previously prepared metal salt solution, and the titration rate was controlled to ensure that the pH of the reaction mixture was about 9. Then, the reaction mixture solution was allowed to react at 70 °C for 12 h under stirring. After the reaction was over, the resulting slurry mixture was filtered and obtained white filter cake. The filter cake was washed with deionized water several times until the pH of the filtrate was approximately equal to 7. Finally, the resulting product was dried at 50 °C under vacuum for 24 h, and then, the product was ground to obtain a powder with particle size of 200 mesh.

Preparation of E-LDHs

Zn (NO3)2⋅6H2O (17.82 g) and Al (NO3)3⋅9H2O (11.25 g) were dissolved in 100 mL of deionized water according to the molar ratio of n(Zn2+): n(Al3+) = 2:1. EDTA (2.92 g) was added to a beaker containing 100 mL of deionized water and stirred until the EDTA was completely dissolved. Subsequently, the previously prepared metal salt solution was slowly added to the stirred EDTA solution. After the addition, the pH of the reaction mixture was adjusted by adding NaOH solution (2 moL/L) until the pH was about 5. After that, the reaction mixture was allowed to continue in a water bath at 70 °C for 5 h under stirring. After the reaction was over, the reaction mixture was naturally cooled to room temperature, filtered and vacuum dried to obtain E-LDHs.

Preparation of AE-LDHs

E-LDHs (1 g), boric acid (0.3 g, catalyst) and aniline solution (50 mL) were added sequentially to a three-necked flask. Then, the reaction mixture was ultrasonically dispersed for 20 min. Subsequently, DMF (5 mL) and toluene were added, and the reaction mixture was allowed to continue react in a water bath at the set temperature for several hours. After completion of the reaction, the reaction mixture was cooled to room temperature and poured into 300 mL of deionized water. After stirring, washed with water (300 mL) ultrasonically and then the reaction mixture was filtered. In order to remove the reactants and by-products, the product was dissolved in 150 mL of ethanol solution (CH3CH2OH) at 55 °C and washed again. Subsequently, the reaction mixture was filtered to obtain the target product. Finally, the resulting product was dried at 50 °C under vacuum. The AE-LDHs modification scheme is shown in Fig. 1.

Fig. 1
figure 1

The synthetic route of AE-LDHs

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Preparation of PLA/AE-LDHs nanocomposites

Poly (lactic acid) (Mw = 100,000 g mol−1) was dried to remove residual water for 24 h at 50 °C under vacuum. PLA/AE-LDHs nanocomposites were prepared by melt blending in a one step process. The weight fraction of AE-LDHs in PLA was 0 wt%, 0.05 wt%, 0.1 wt%, 0.3 wt%, and 0.5 wt%, respectively. PLA/AE-LDHs nanocomposites were named as PLA0, PLA1, PLA2, PLA3, PLA4, respectively. They were extruded by double-screw extruder (HAAKE Mini-Lab II, Thermo Fisher Scientific, Karlsruhe, Germany, diameter 19 mm and L/D = 25:1) at 180 °C, and the screw speed was set at 30 rpm. The products were dried in a desiccator for 24 h. Subsequently, these products were extruded, pelletized and injection-molded (HAAKE Mini Jet Pro, Thermo Fisher Scientific) into standard test specimens for tensile tests, impact tests, and dynamic mechanical tests. During injection molding, all samples were molded under 80 MPa pressure at 190 °C, mold temperature of 65 °C, and pressure-holding time of 30 s.

Characterization

Fourier transform infrared spectroscopy

Fourier transform infrared (FTIR) spectra of the samples were recorded on a Bruker Equinox 55 (Karlsruhe, Germany) spectrometer using standard KBr pellet/disk technique. Spectra were recorded over a spectral range of 4000 to 400 cm−1, with a resolution of 4 cm −1. All the samples were dried in a vacuum oven at 60 °C for 24 h before analysis.

XPS analysis

The chemical compositions of E-LDHs and AE-LDHs were determined by X-ray photoelectron spectrometer (XPS) using an instrument, Axis Ultra (Kratos Analytical, Britain) that employed a monochromatic Al Kα X-ray source (284.7 eV). All binding energies were referenced to the C1s hydrocarbon peak at 284.7 eV.

X-ray diffraction

X-ray diffraction (XRD) patterns were obtained using a Philips X’Pert X-ray diffractometer (Almelo, the Netherlands) with Cu Kα X-ray source at λ = 0.1540 nm (50 kV, 35 mA). Diffraction spectra were obtained over a 2θ range of 5–50° under a scanning speed of 0.10 °/s.

Static contact angle test

Static contact angles of water on LDHs, E-LDHs and AE-LDHs surfaces were determined using contact angle meter (JJ2000B2, Zhongchen Digital Technology Co., China). The samples were pressed into small sheets of thickness of about 2 mm by a SSP-10 type tablet press before testing.

Transmission electron microscopy

The morphology of LDHs, E-LDHs and AE-LDHs were observed by a transmission electron microscopy (TEM), TECNAI G2-F20 (FEI Co., the Netherlands). Samples were observed under an acceleration voltage of 75 kV.

Scanning electron microscopy

Scanning electron microscopy (SEM) (Inspect F, FEI Instrument Co., Ltd., the Netherlands) at 20 kV of accelerating voltage was applied to observe the impact fracture morphology of PLA/AE-LDHs nanocomposites. Before SEM analysis, the samples were sputter coated with gold to limit charging during the tests.

Thermogravimetric analysis

Thermogravimetric analysis (TGA) was conducted on SDT-Q600 thermogravimetric analyzer (USA). Each sample (about 5–10 mg) was heated at a heating rate of 10 °C min−1 in the temperature range of room temperature to 700 °C under nitrogen atmosphere (100 mL min−1).

Differential scanning calorimetric

Differential scanning calorimetry (DSC) was conducted on a DSC Q2000 differential scanning calorimeter (USA). Each sample was heated from 20 °C to 180 °C at a heating rate of 10 °C min−1 under nitrogen atmosphere (50 mL min−1).

Polarized optical microscopy

The spherulitic morphologies of the samples were studied by polarized optical microscopy (POM) (Linkam-CSS450, England). The samples were melted at 200 °C for 2 min and then they were rapidly cooled to 130 °C for isothermal crystallization. Micrographs of the isothermally crystallized spherulites of PLA and PLA nanocomposites were studied.

Dynamic mechanical analysis

Dynamic mechanical analysis (DMA) was operated in the tensile mode on a DMAQ800 analyzer (USA). Temperature scans at 1 Hz frequency were carried out from room temperature to 120 °C at a heating rate of 3 °C min−1. The samples were acquired from injection molding and dried in a vacuum oven at 40 °C for 24 h before testing.

Tensile testing

Tensile tests were carried out according to ASTM D 638 using Instron 4302 Universal Testing Machine (Norwood, MA, USA) at a cross-head speed of 5 mm min−1. At least five measurements for every sample were recorded. Samples of PLA nanocomposites were standard test specimens from the injection mold.

Notched Lzod impact testing

Notched Lzod impact testing of the samples were carried out using a Testing Machines Inc. (TMI) 43–02-01 Monitor/Impact machine (New Castle, DE, USA) according to ASTM D256. Two-millimeter-deep notches were cut into specimen beams using a TMI notch cutter. All results presented were the average values of five measurements.

Rotational rheological measurement

The rheological behaviors of all samples were studied using a stress-controlled rotational rheometer (AR 2000 ex, TA Instruments) with 25 mm diameter and 1 mm gap parallel plates. Prior to the measurement, each disk-shaped sample was prepared by compression molding, which was heated to 170 °C and held for 3 min to completely melt PLA along with the nucleating agents. Deformation in the linear viscoelastic region was observed to be 2% and chosen for all subsequent frequency swept tests. Samples were equilibrated at 170 °C before testing. The frequency was varied from 0.1 to 100 rad/s.

Capillary rheological analysis

The capillary rheological analysis of the samples was tested by trace a mixed rheometer (MiniLab II, Thermo Fisher Scientific). The samples of 7 g were melted at 180 °C for 10 min and then were sheared from 1 to 50 rpm at 180 °C.

Results and discussion

Design of orthogonal experiments

Orthogonally designed experimental methods, were used to choose the optimum reaction conditions for the synthesis of AE-LDHs by amidation. The factors for the orthogonal experimental included mass ratios of E-LDHs and boric acid (H3BO3), reaction temperature, reaction time, which are presented in Table 1 [24]. The scheme for orthogonal experiments and effects of the modification are presented in Table 2. Besides, range analysis was performed choosing the interfacial contact angle values as the objective function (Table 2). The analytic results are presented in Table 2, in which the impact factors of each factor range (R) for amide modification of LDHs were 10.67, 9.00, and 3.67, respectively. From these results, it was clear that temperature had the greatest impact on the modification of LDHs. The results of orthogonal experiments showed that the optimum reaction conditions were a 1:0.3 mass ratio of E-LDHs and H3BO3, reaction temperature of 70 °C and reaction time of 2 h.

Table 1 The orthogonal experiment design of the synthesis for AE-LDHs
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Table 2 The orthogonal results of the AE-LDHs synthesis
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Hydrophilicity or hydrophobicity of material surface became highly obvious from the static contact angles of the water on the material surfaces and comparison of their wetting behaviors. Images obtained during the interface contact angle measurements are shown in Fig. 2. The water contact angles of LDHs, E-LDHs and AE-LDHs were 32°, 51°, and 87°, respectively. Compared with LDHs and E-LDHs, the water contact angle of the AE-LDHs was increased significantly. This result showed the improvement in hydrophobicity of AE-LDHs after modification by amidation reaction, which was beneficial for improving its compatibility with PLA [25]. This was due to the fact that aniline groups were grafted onto the E-LDHs by amidation, which improved the hydrophobicity of AE-LDHs.

Fig. 2
figure 2

Contact angels of LDHs, E-LDHs, AE-LDHs: a contact angle of LDHs; b contact angle of E-LDHs; c contact angle of AE-LDHs

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FTIR analysis

Figure 3a shows the FTIR spectra of LDHs, E-LDHs, and AE-LDHs. It could be seen that the spectrum of LDHs showed a strong and broad absorption peak at about 3460 cm−1, due to the stretching vibrations of the O-H group. The band observed around 1638 cm−1 was assigned to the bending vibration of interlayer water molecules. The absorption peak at 1385 cm−1 was the characteristic telescopic vibration peak of NO3−, which indicated that the LDHs were successfully prepared. Feng et al. [26] also reported similar results. In the E-LDHs spectrum, the absorption peaks at 1608 cm−1 and 1169 cm−1, due to the stretching vibration of the easter bond and C-N, respectively. In addition, disappearance of the characteristic peak at 1385 cm−1 indicated the complete replacement of the NO3− groups. This proved that the E-LDHs was successfully prepared. In the AE-LDHs spectrum, the absorption peaks at 3103 cm−1 and 698 cm−1, due to the N-H stretching and bending vibrations, respectively. The amide I peak at 1650 cm−1 was due to the C=O stretching vibration in amide groups. The amide II peak at 1415 cm−1 could be attributed to the bending vibrations of C-N-H. The amide III peak at 1258 cm−1 could be attributed to the combination of C-N stretching and N-H bending vibrations. The absorption peak at 766 cm−1 was due to the out-of-plane bending vibrations of aromatic C-H [27, 28]. Results of FTIR indicated that the AE-LDHs were successfully synthesized. FTIR analysis of PLA0 and PLA3 are shown in Fig. 3b. In the melting processing, due to the shear effect of the rotor, which led to the intra-and inter-molecular hydrogen bonds of PLA were broken. Once PLA molecular chain was destroyed, it would produce a large number of free hydroxyl groups. It was expected that the free hydroxyl groups could form a new hydrogen bond between AE-LDHs molecules and PLA molecules during the extrusion process, resulting in intermolecular complexation and plasticization. As shown in Fig. 3b, the spectrum of PLA0 showed an absorption peak at about 3505 cm−1, due to the stretching vibrations of the O-H group; the absorption peak at 1757 cm−1, corresponding to the C=O characteristic absorption peak. The O-H stretching vibration peaks of PLA/AE-LDHs nanocomposites shifted to a lower wave number, which indicated that the hydrogen bonds could be re-formed between O-H groups in the PLA and N-H groups in the AE-LDHs since the interactions between O-H groups were weaker in the PLA/AE-LDHs nanocomposites than that in the neat PLA. When the AE-LDHs nucleating agent was added, the C=O characteristic absorption peak was obviously shifted to a low wave number of 1753 cm−1. This phenomenon could be explained by the fact that the N-H in the AE-LDHs and the C=O in the PLA formed hydrogen bonds, causing the charge to separate from the C=O bonds into other atoms. The double bond property of C=O was reduced,which led to the characteristic peak of C=O shifted toward a low wave number [29]. Moreover, the movement of these peaks confirmed that the PLA and the AE-LDHs were not simply mixed, but through an intermolecular interaction to form an ordered structure.

Fig. 3
figure 3

FT-IR characterization: a FT-IR spectra of LDHs, E-LDHs, AE-LDHs; b FT-IR spectra of PLA0 and PLA3

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XPS analysis

XPS was applied to investigate the changes in functional groups and elemental compositions of E-LDHs, and AE-LDHs. The XPS spectra of E-LDHs and AE-LDHs are shown in Fig. 4. The C/O atomic ratio (34.38%/60.59%) of E-LDHs confirmed that organic intercalated LDHs was successfully prepared. The C1s peaks in the XPS spectra of AE-LDHs were divided into five peaks, which could be ascribed to carbon atoms with different functional groups, corresponding to C-C (284.6 eV), C-O (286.1 eV), C-O-C (287.2 eV), C=O (288.5 eV), and O=C-O (292.5 eV) bonds [30]. It can be seen from Table 3 that the nitrogen content of AE-LDHs increased to 10.42% compared with 5.03% of E-LDHs. The atomic ratio of N/C for E-LDHs was 0.146, which increased to 0.171 for AE-LDHs. Moreover, the ratio of O/C decreased from 1.76 for E-LDHs to 0.47 for AE-LDHs. This proved that the amidation reaction was successful [31]. This result was consistent with the FTIR analysis.

Fig. 4
figure 4

The XPS analysis of E-LDHs and AE-LDHs: a full spectrum of E-LDHs; b C spectrum of E-LDHs; c N spectrum of E-LDHs; d full spectrum of AE-LDHs; e C spectrum of AE-LDHs; f N spectrum of AE-LDHs

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Table 3 XPS analysis of E-LDHs and AE-LDHs
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XRD analysis

XRD can be used to observe the crystalline structure of materials. Figure 5a and b show the XRD patterns of LDHs, E-LDHs, AE-LDHs, PLA0 and PLA3. In Fig. 5a, the characteristic peaks of LDHs appeared at 2θ = 9.9° (first peak), 19.9° and 34.4°, respectively. According to Bragg’s equation (ΔHm, n = 1), corresponding to interlayer spacing was 0.89 nm, 0.45 nm and 0.14 nm, respectively. The diffraction peaks of LDHs were sharp with high peak intensity, which suggested that the LDHs crystal was regular, and the crystallinity and degree of orientation were high. This result was consistent with that reported by Chai et al. [32]. The diffraction peaks of E-LDHs tended towards the lower angle regions, in which the first characteristic peak of E-LDHs was appeared at 2θ = 7.4° and the layer spacing was increased to 1.20 nm. This result indicated that EDTA was successfully intercalated into LDHs, resulting in an increase in interlayer spacing. Moreover, the first characteristic peak of AE-LDHs was appeared at 2θ = 5.8°, and the interlayer spacing was increased to 1.52 nm. This was mainly attributed to the introduction of amide bonds in LDHs, and the interlayer distance was further increased. As shown in Fig. 5b, the characteristic peaks of PLA appeared at 2θ = 16.74°, 19.13°, and 22.51°, respectively. This indicated that the lattice structure of PLA showed an alpha crystal shape. For the PLA/AE-LDHs nanocomposites, the similar characteristic peaks were observed, which suggested that the incorporation of AE-LDHs did not change the crystal structure of PLA. Moreover, compared with neat PLA, the diffraction peaks of PLA nanocomposites became sharp and neat, which indicated that the introduction of AE-LDHs could significantly improve the crystallization of PLA. Moreover, in Fig. 5a, the diffraction peak of AE-LDHs was observed at 2θ = 5.8°, while it was disappeared in the PLA3, suggesting that the layered AE-LDHs had been exfoliated in the nanocomposites. Ahmed et al. also reported similar results [33]. These results were in good agreement with the DSC measurements.

Fig. 5
figure 5

X-ray diffraction spectra: a X-ray diffraction spectra of LDHs, E-LDHs, AE-LDHs; b X-ray diffraction spectra PLA0 and PLA3

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TEM analysis

In order to further clarify the macrostructures of LDHs, E-LDHs, and AE-LDHs. The samples were first processed by embedding method (The embedding agents were made of Epon 812, MNA, DDSA, and DMP-30 epoxy resins) then placed on a microgrid to prepare ultra thin slices and studied by transmission electron microscopy. Figure 6a, b, and c show the TEM images of LDHs, E-LDHs, and AE-LDHs at the same magnification. It was clear that both LDHs and E-LDHs presented layered structures, but the morphology of LDHs showed plate-like layered with smooth surfaces. Moreover, the thickness of E-LDHs sheet was larger than that of LDHs because of the intercalation of EDTA anions into LDHs layers. Compared with E-LDHs, the thickness of AE-LDH increases and the lamellar becomes irregular, indicating that the amide bond was successfully grafted onto E-LDHs and could thus slightly affect the crystal structure of AE-LDHs, which was in accordance with the FTIR, XPS, and XRD results [34]. Figure 6d shows the TEM image of the formation of PLA nanocomposites with 0.3 wt% AE-LDHs content. The dark regions represent the AE-LDHs layers, whereas the gray regions correspond to the PLA matrix. it could be seen that most of the AE-LDHs nano-layers were exfoliated, implying a nano-size level of dispersion. Moreover, the AE-LDHs layers were evenly dispersed in the PLA matrix, indicating that the nanocomposite had the intercalated and partially exfoliated structures [35]. Combining XRD with TEM results proves PLA3 as a nanocomposite.

Fig. 6
figure 6

TEM images of LDHs, E-LDHs, AE-LDHs, and PLA/AE-LDHs nanocomposites: a LDHs; b E-LDHs; c AE-LDHs; d PLA3

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Mechanical properties

For polymeric materials, their mechanical properties are so critical to practical processing applications. Therefore, it is necessary to improve the mechanical properties of polymeric materials. Mechanical properties are mainly manifested in three aspects of maximum tensile strength, elongation at break and impact strength. The effects of varying the amount of AE-LDHs on the mechanical properties of neat PLA are shown in Fig. 7. The tensile strength of pure PLA was 80.76 MPa, elongation at break of pure PLA was 18.47%, and impact strength of pure PLA was 10.30 KJ/m2. It can be seen from the Fig. 7 that the addition of AE-LDHs significantly improved the mechanical properties of the PLA nanocomposites. When the AE-LDHs dosages reached 0.3 wt%, the mechanical properties of PLA nanocomposites reached maximum values. This composite presented 84.15 MPa of tensile strength, 22.12% of elongation at break, and 19.81 KJ/m2 impact strength, an increase of 4.68%, 22.17%, and 92.91%, respectively. However, when the AE-LDHs content was above 0.3 wt%, the mechanical performance of PLA/AE-LDHs nanocomposites decreased gradually. This can be explained by the fact that the incorporation of a small amount of AE-LDHs led to an increase in the crystal nucleus and reduction in its grain size [36]. On the other hand, the addition of the nucleating agent promoted the crystallization of PLA, structured the arrangement of macromolecular chains, enhanced the intermolecular forces, and eventually increased the mechanical properties of the material. Furthermore, a strong interaction between the PLA molecular chains and exfoliated AE-LDHs sheets increased their compatibility at the interface, and consequently the mechanical performance was improved. However, with increase in the amount of AE-LDHs, the tensile strength, elongation at break, and impact strength of the PLA nanocomposites gradually decreased. The main reasons are as follows: on one hand, with increase in the amount of AE-LDHs, the number of groups on PLA that could interact with AE-LDHs decreased. This led to excessive accumulation and formation of agglomerates between the biphasic interface through strong electrostatic interactions, which greatly restricted the movement of PLA molecular chains and the growth space for its chains and destroyed the continuity of the matrix in the composite. On the other hand, the excessive AE-LDHs particles accumulate with each other, resulting in an increase in their size, which led to form accumulate stress in PLA and caused a reduction in mechanical properties [37, 38].

Fig. 7
figure 7

Effects of AE-LDHs on mechanical properties of PLA/AE-LDHs nanocomposites: a Tensile strength and elongation at break; b Impact strength and fracture scanning electron microscopy morphology of PLA0 and PLA3

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SEM can be employed to investigate the fracture characteristics of PLA and PLA nanocomposites, which confirmed the effects of addition of AE-LDHs nucleating agents on the mechanical properties of the materials. Figure 7b shows the morphologies of impact fractured surfaces of pure PLA and PLA nanocomposites. The PLA3 with the best mechanical performance was selected as the sample compared with pure PLA. It can be seen that the fractured surface of pure PLA was relatively flat and parallel lines appeared, and its fractured surface without occurred plastic deformation, which showed a typical brittle fracture. Brittle fracture leads to a low toughness and shock strength, resulting in poor mechanical properties of PLA. In Fig. 7b, the AE-LDHs layer was coated with PLA matrix, which was dispersed in the PLA matrix in a disorderly manner. However, the roughness of the fractured surface of PLA3 was obviously increased, and there was a significant scaly epitaxial deformation compared to pure PLA. In addition, a large number of oriented fibers and tough cavities were observed, which apparently showed characteristics of toughness fracture and improved its toughness on the incorporation of AE-LDHs. This could be explained by the fact that AE-LDHs nucleation agents could induce energy dissipation mechanisms in PLA, which could decelerate the initiation and propagation of the crack [39]. So, the toughness of PLA was improved. These results indicated that the crystallization of PLA was promoted, which was in good agreement with the results from the analysis of mechanical properties.

Thermogravimetric analysis

In order to investigate the effect of AE-LDHs on the thermal stability of PLA matrix. TGA-DTG analysis was employed to study the thermal degradation of AE-LDHs, pure PLA and PLA/AE-LDHs nanocomposites The thermal degradation curves of AE-LDHs, pure PLA and its nanocomposites under nitrogen flow at a heating rate of 10 °C min−1 are depicted in Fig. 8, which shows three decomposition stages. The first decomposition stage below 100 °C can be attributed to the loss of physically adsorbed water. The second decomposition stage between 100 and 360 °C can be attributed the thermal decomposition of the organic groups. Consequently, a small decomposition peak appears in the DTG curve. Furthermore, the third decomposition stage from 360 to 600 °C corresponds to the combustion of the residue, and the DTG curve demonstrated one strong decomposition peak at 376.6 °C [40]. Figure 8 also exhibits the initiation temperature (Tonset), maximum temperature (Tmax), and final decomposition temperature (Tend) of the samples. The Tonset, Tmax, and Tend for PLA0 was 290.8 °C, 359.7 °C, 374.1 °C. Compared with pure PLA, PLA3 showed 299.1 °C of Tonset, 365.4 °C of Tmax, and 381.9 °C of Tend, an increase of 8.3 °C, 5.7 °C, and 7.8 °C, respectively, owing to the incorporation of AE-LDHs. These results showed significant improvement in the thermal stability of PLA nanocomposites. This could be explained as follows: on one hand, AE-LDHs were uniformly dispersed in the PLA matrix. Therefore, the interaction between AE-LDHs and PLA was relatively strong, which restricted the movement of PLA molecular chains. Due to the fact that PLA molecular chain motions were inhibited, which increased the energy barrier for thermal decomposition of PLA, resulting in heat resistance was improved. On the other hand, due to the superior nucleation effect of AE-LDHs, the crystallization behavior of PLA was improved, and the crystallinity was increased, thereby improving the thermal stability of PLA nanocomposites [41, 42]. Moreover, compared with pure PLA, it should be noted that the final weight of PLA nanocomposites was higher than PLA0 because of the presence of the AE-LDHs char residue [43]. The increased char residue could obstruct heat and mass transfer path, further improved the thermal stability of composite.

Fig. 8
figure 8

Thermogravimetric curves for AE-LDHs, PLA0 and PLA3: a thermogravimetric curves for AE-LDHs, PLA0 and PLA3; b differential thermogravimetric curves for AE-LDHs, PLA0 and PLA3

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Differential scanning calorimetric analysis

Thermal properties of PLA and PLA nanocomposites were studied with differential scanning calorimeter analysis. In general, the crystallization process was divided into heating process and cooling process. In this study, the non-isothermal cold crystallization process of pure PLA and its nanocomposites was further studied with DSC in a wide range of crystallization temperatures from 20 °C to 180 °C [44]. Figure 9a shows the DSC heating curves of PLA0 and PLA3. The degree of crystallinity (Xc) was determined from DSC analysis data employing Eq. 1, where ΔHm, ΔHc were the measured melting enthalpy, cold-crystallization enthalpy of PLA, respectively. The degree of crystallinity (Xc) of the polymer could be calculated by using the equation given below:

$$ {X}_{\mathrm{c}}=\frac{\Delta {H}_{\mathrm{m}}\hbox{-} \Delta {H}_{\mathrm{c}}}{f\times \Delta {H}_{\mathrm{m}}^o}\times 100\% $$
(1)

where \( \Delta {H}_{\mathrm{m}}^o \) was the melting enthalpy expected for a polymer with 100% crystallinity [45]. The heat of fusion \( \Delta {H}_{\mathrm{m}}^o \) of pure PLA was 93.7 J/g [46], and f was the weight fraction of pure PLA. In Fig. 9, Tm and Tc represented the measured melting temperature and cold-crystallization temperature, respectively. It can be seen from Fig. 9a that the values of Tm and Tg tended towards the lower temperature regions. The value of Tm of PLA0 was 169.50 °C, whereas the Tm of PLA3 was 164.40 °C. This result suggested that the formation of hydrogen bonds and the breaking of intermolecular interactions in PLA after introduction of AE-LDHs. Moreover, the cold crystallization peak of PLA nanocomposites disappeared after the introduction of AE-LDHs. This was due to the addition of AE-LDHs reduced the activation energy of the PLA molecular chain rearrangement, which made the crystallization easier and controlled the cold crystallization behavior of PLA. Moreover, according to the data of Fig. 9a could be calculated that the Xc of PLA0 to be 5.21%, whereas the Xc of PLA3 increased to 35.36%, which also demonstrated that the AE-LDHs improved the crystallinity of PLA. Figure 9b exhibits the DSC cooling curves for PLA0 and PLA3 and also presents the thermal parameters of PLA0 and PLA3 from the melt at a cooling rate of 2 °C min−1. As shown in Fig. 9b, addition of AE-LDHs increased the crystallization peak temperature (Tmo) and initial crystallization temperature (To) of PLA3 in the cooling crystallization process, compared with that of neat PLA. The crystallization enthalpy of the PLA nanocomposites increased from 13.48 J/g to 40.19 J/g, which was an increase of 198%. This can be attributed to AE-LDHs had good interfacial compatibility with PLA, could be homogeneously dispersed, and AE-LDHs played the role of a heterogeneous nucleating agent during the crystallization process that improved the crystallization of PLA [47, 48]. These results further confirmed that AE-LDHs could be used as the heterogeneous nucleating agent to promote the crystallization of PLA, which were also in good accordance with the mechanical testing results.

Fig. 9
figure 9

Differential scanning calorimetry curves of PLA0 and PLA3: a heating stage; b cooling stage

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Polarized optical microscopic analysis

It is of great interest to study the isothermal crystallization behavior to obtain information useful for the industrial applications of PLA. The crystallization process of a polymer mainly includes nucleation and crystal nucleus growth. So, the crystallization rate mainly includes nucleation rate, crystal growth rate, and total crystallization rate. The crystallization behavior of PLA0 and PLA3 were investigated by POM. Figure 10 shows the POM images of isothermal crystallization process of PLA0 and PLA3 at 130 °C for 20 min. As expected, PLA exhibited a typical spherulitic morphology. After 5 min of isothermal crystallization, the spherulites of PLA showed a distinct black extinction phenomenon. It was clear that the spherulite number of PLA3 was much larger than that of PLA0 under the same conditions. This could be attributed to the even distribution of the nucleating agent, which led to an increase in the nucleation density of PLA. Figure 11 presents the spherulite radius of PLA0 and PLA3, which changes with time during the isothermal crystallization at 130 °C. It can be seen from the figure that the spherulitic radius of PLA0 and PLA3 increases linearly with time. Hence, the growth rate (G) of the spherulites can be observed from the slope of two straight lines in the figure. By fitting, the spherulite growth rates of PLA0 and PLA3 were 3.04 μm/min and 1.17 μm/min, respectively. This was due to the impingement of large number of spherulites with each other, which severely restricted the growth of spherulites [49]. In conclusion, the total crystallization rate of PLA3 was much higher than that of pure PLA. Results of POM analysis showed that AE-LDHs could be used as a novel nucleating agent to improve the crystallization rate of PLA, which was consistent with the aforementioned DSC results.

Fig. 10
figure 10

Time evolution of polarized optical microscopy morphology of PLA0 and PLA3: a PLA0 (5 min); b PLA0 (10 min); c PLA0 (20 min); d PLA3 (5 min); e PLA3 (10 min); f PLA3 (20 min)

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Fig. 11
figure 11

Changes of spherulite radius as functions of time PLA0 and PLA3 during isothermal crystallization at 130 °C

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Dynamic mechanical analysis

The dynamic viscoelasticity of the polymer during stretching is usually analyzed using a dynamic thermo mechanical instrument. The storage modulus represents the rigidity of a viscoelastic material and is proportional to the maximum energy stored during a loading cycle [50]. Storage modulus can be used to reflect the elasticity of the viscoelastic material. The more the storage modulus is, the greater the rigidity of the material is, and the more difficult the material deformation is. The loss modulus represents the energy dissipated due to viscous deformation when the material is deformed. The loss tangent (tan δ) is defined as the ratio of loss modulus to storage modulus. It is a measure of the energy lost, which is used to characterize the loss damping or internal friction in a viscoelastic system. Variations of storage modulus and tan δ of the samples with temperature are shown in Fig. 12a and b. It can be seen from Fig. 12a that the curves of PLA0 and PLA3 can be divided into three regions: glassy region (30–50 °C), transition region (50–60 °C), and rubbery region (60–100 °C). In the glassy region, the motions of molecular chains of the samples were frozen and they went through a glassy state, and the storage modulus basically remained constant in the higher range. In the glassy transition region, there was a large drop of the storage modulus of PLA0 and PLA3. This can be explained by the fact that the molecular chains in this region could move freely and lead to a rapid drop of the storage modulus. However, the storage modulus of the PLA3 was lower than that of pure PLA, which could be due to the influence of AE-LDHs on the mechanical performance of PLA molecular chains. As shown in Fig. 12b, the peak of the loss factor was the glass transition temperature (Tg) of the samples. The Tg of PLA0 and PLA3 were 70.64 °C and 57.75 °C, respectively. Such a drop in Tg of PLA3 was due to the fact that the improvement of crystallinity of PLA/AE-LDHs nanocomposites increased flexibility of PLA chain, and reduced Tg of PLA. Hence, the glass transition temperature (Tg) and storage modulus of PLA nanocomposites decreased. These results showed that the incorporation of AE-LDHs could effectively reduce the damping factor in the glass transition region of PLA, and the elasticity and compatibility of the PLA were improved [51, 52]. The DMA results were in good accordance with the DSC results.

Fig. 12
figure 12

Dynamic mechanical relaxation behavior of PLA0 and PLA3: a storage modulus; b tan δ

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Rotational rheological analysis

Figure 13 exhibits the storage modulus (G’), loss modulus (G”) and complex viscosity (η) as a function for the PLA and its nanocomposites. The shear frequency region is divided into three parts: low frequency region (0–20 Hz), intermediate frequency region (20–60 Hz), high frequency region (60–100 Hz). As shown in Fig. 13a and b, the storage modulus and loss modulus of the sample increased with increase of angular frequency (w), which showed to a typical linear rheological behavior. It was worth noting that the storage modulus and loss modulus of all samples had a turn in the intermediate frequency region, which was due to the loose shape of the dispersed phase of their system and the presence of its phase interface. It was clear that the storage modulus and loss modulus of the PLA/AE-LDHs (0.3 wt%) nanocomposites were increased obviously compared with pure PLA, in the given frequency range. This was because the addition of AE-LDHs acts as a cross-linking point, which restricted the movement of the molecular chains, and the interaction between AE-LDHs and PLA macromolecular caused the high flow resistance. In this case, the degree of entanglement between the molecular chains increases, so that the storage modulus and loss modulus of PLA nanocomposites increase. Dumitriu et al. [53] reported similar behavior of increase in melt state rheological properties of polyester. However, when the AE-LDHs content was above 0.3 wt%, the values of G’ and G” of PLA nanocomposites decreased slightly, but still higher than that of PLA0. This was consistent with the aforementioned mechanical properties results. Figure 13c showed the relationship between complex viscosity and angular frequency of the PLA/AE-LDHs nanocomposites. As shown in Fig. 13c, PLA0 and PLA nanocomposites showed transitions to non-Newtonian pseudoplastic behaviors, due to the viscosities decreased consistently with increase in frequency in the deformation region of 0.01–100 Hz. This revealed a typical feature, in which the frequency was completely dependent on the nature of the sample [54]. This can be explained by the fact that as the effect of shear thinning flow increased, the flow resistance decreased. Moreover, the complex viscosity of PLA3 was significantly higher than that of pure PLA. This could be that the amide bonds of AE-LDHs and the ester bonds of PLA form hydrogen bonds, which increased the binding force inside the material. In conclusion, the interfacial interactions between the PLA chains and nucleating agents on their surface increased the storage modulus, loss modulus and complex viscosity of PLA nanocomposites.

Fig. 13
figure 13

Rheological behavior of PLA0–4: a storage modulus; b loss modulus; c complex viscosity

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Capillary rheological analysis

Figure 14a and b show the capillary rheological curves of PLA0 and PLA3. It could be seen from Fig. 14a that the apparent viscosity (ln () of PLA0 and PLA3 showed similar decreasing trends, with an increase of shear rate at constant temperature, which indicated that the melt was pseudo-plastic fluid with shear-thinning phenomena. It is well known that shear thinning phenomena often occurs in polymer melt processing, which is related to the molecular structure of the material. However, it can be seen from Fig. 14a that the apparent viscosity ln η for PLA3 was obviously higher (by 85.86%) than that of PLA0 under the same shear rate. This was due to the fact that long and flexible macromolecular chains of the PLA/AE-LDHs nanocomposite could entangle each other, which led to a decrease in the liquefaction of the composite material, and the ln η value increased [55, 56]. It can be seen from Fig. 14b that the torque of PLA and its composite increased rapidly within a short period of time and then decreased gradually. This could be attributed to the material was not completely melted at the beginning of the feed process, resulting in a gradually increased in resistance in the mixing chamber, and so the torque increased. Subsequently, after the feeding process was finished, the material was compacted and began to melt under the shear of the rotor. With further melting of the material, the resistance was gradually decreased, resulting in the gradual reduction of torque. Furthermore, It is evident from Fig. 14b that the value of maximum torque for PLA3 was greater (by 46.21%) than that of pure PLA. This was due to the fact that the formation of hydrogen bonds between AE-LDHs and PLA led to an increase in intermolecular forces, resulting in a decrease in the fluidity of the PLA composite material and an increase in torque. The result was in good accordance with the result shown in Fig. 14a.

Fig. 14
figure 14

Capillary rheological of PLA0 and PLA3: a ln η-shearing rate curves of PLA0 and PLA3; b torque-time curves of PLA0 and PLA3

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Conclusions

In this paper, a new nucleating agent (AE-LDHs) was prepared by the amidation of E-LDHs. The conditions of synthetic AE-LDHs were optimized by orthogonal experiment. The structure and surface properties of LDHs and AE-LDHs were analyzed by FTIR, XPS, XRD, and static contact angle measurements. Subsequently, PLA/AE-LDHs nanocomposites were prepared by incorporation of AE-LDHs by melt bending and tested for its performance. Mechanical tests showed that the mechanical properties of PLA/AE-LDHs nanocomposites were significantly enhanced compared with pure PLA. DSC and POM results indicated that the crystallinity and crystallization rate of PLA/AE-LDHs nanocomposites were improved significantly. Rheological analysis demonstrated that the storage modulus, and loss modulus increased to a great extent and the viscosity of the blend decreased with increasing frequency. Hence, numerous test results showed that AE-LDHs could be used as a high-efficiency nucleating agent to modify PLA and provide a large number of heterogeneous nucleation sites to improve the crystallization rate and crystallization properties of PLA.