Xiaolong Jiaab,
Tianhong Tanga,
Dan Chenga,
Lijuan Guoa,
Cuihua Zhanga,
Qing Cai*a and
Xiaoping Yangab
aState Key Laboratory of Organic-Inorganic Composites, College of Materials Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, P. R. China. E-mail: caiqing@mail.buct.edu.cn; Fax: +86 10 64412884
bChangzhou Institute of Advanced Materials, Beijing University of Chemical Technology, Jiangsu 213164, P. R. China
First published on 19th November 2014
Polyacrylonitrile (PAN) electrospinning in combination with sol–gel method has been a common technique to produce inorganic nanoparticles containing composite carbon nanofibers (CNFs) for diverse applications. To investigate the morphology evolution and crystal transformation of inorganic components along with CNF formation, bioactive glass (BG) containing CNFs (CNF/BG) were prepared by sintering as-spun PAN/precursor composite nanofibers in a nitrogen atmosphere at temperatures of 800, 1000 and 1200 °C. Comprehensive characterizations were performed with TEM, SEM-EDXA and XRD. For samples sintered at 800 °C, numerous BG nanoparticles were observed inside the CNFs and mainly in an amorphous state. With the sintering temperature raised to 1000 °C, a number of spherical BG nanoparticles were detected on the surface of the resulting CNFs, with a crystal structure of wollastonite (β-CaSiO3) polycrystals. When the samples were sintered at 1200 °C, the BG nanoparticles on the surface of CNFs merged into forms with cuboid-like geometry, mainly consisting of pseudowollastonite (Ca3(Si3O9)) single crystals. Based on the geometry evolution and dynamic size distribution function analyses (Ostwald ripening and Smoluchowski equations), it was concluded that the growth of BG nanoparticles conformed to the ripening mechanism at 800 °C and migration–coalescence mechanism at 1200 °C, while the process involved both ripening and migration–coalescence mechanisms at 1000 °C.
Noticeably, physicochemical properties and surface morphology of CNF/M or CNF/MOx composites were found significantly depending on preparation parameters such as pre-oxidation conditions and sintering conditions, etc.13–15,20,24 Wang et al. reported that the size of Fe3O4 nanoparticles in CNF/Fe3O4 composite increased remarkably along with higher sintering temperature. They found the electrical properties of CNF/Fe3O4 films were optimized for samples sintered at 600 °C.20 Our previous works revealed that Sn nanoparticles had strong tendency to migrate to the surface of CNFs when the sintering temperature in producing CNF/Sn composites was raised from 550 to 1000 °C. At the same time, the initial amorphous Sn transformed into crystalline Sn. These transitions led to CNF/Sn composites sintered at 850 °C showing the best electrical performance.24 The variances in publications distinctly demonstrated the complexity in preparing such CNF composites and in relating the specific features of M (or MOx) nanoparticles to the final performance. There was a large uncertainty of understanding the morphology evolution and formation mechanism of M (or MOx) nanoparticles during the preoxidation and carbonization of PAN/precursor nanofibers at different sintering temperatures and times. To design and fabricate high performance CNF/M or CNF/MOx composites, thus, it was necessary to investigate the effective approach to control the structure, morphology and size of M (or MOx) nanoparticles in CNFs, which would help to provide theoretical foundation.
Several mechanisms (e.g. Ostwald ripening (OR) and particle coalescence) have been discussed in the context of particle growth dependent on the dominating process of material transport.25–27 It was stated that mass transport was the dominant process to control the particle growth process, although the driving force was associated with the decrease in chemical potential due to the reduction of interfacial area. The difference between Ostwald ripening and particle coalescence mechanism was in their mass transport process. In particular, Ostwald ripening was caused by the diffusion of atoms through the matrix leading to the growth of larger particles at the expense of smaller ones, and atomic diffusion could occur in solid embedding media. While coalescence mechanism involved the collision of migrating particles and subsequent recrystallization and reshaping of aggregates, in which, particle migration required a mobile environment or self-mobile ability. Thus, diffusion-driven aggregation and coalescence were expected only at temperatures above the softening temperature of the embedding medium or the melting temperature of the particles. These mechanisms provided feasible entry paths to look into the evolution of M (or MOx) nanoparticles in preparing CNF/M or CNF/MOx composites.
In this study, bioactive glass (BG) containing CNFs (CNF/BG) was chosen as the model system for the purpose to investigate the morphology evolution and crystal transformation of inorganic components along with CNF formation. As bioceramic material, BGs have been long investigated for applications in bone regeneration.28–30 They were bioactive and able to induce apatite formation by releasing sodium, calcium and silicium ions etc., when they were in contact with body fluid.28,29 In addition to their excellent bone bonding properties, soft connective tissues were also reported able to form a bond with BGs.30 In our previous work,31 CNF/BG composites containing BG nanoparticles of different compositions have demonstrated strong stimulations in biomineralization and osteocompatibility in comparison to pure CNFs. It interested us to investigate the morphology evolution, crystallization behavior and growth kinetics of CNF/BG composites in detail, which was essential to develop new biomaterials for orthopedic applications. Additionally, the study could also provide understanding the growth mechanism of M/MOx nanoparticles in CNF composites prepared in similar procedures. Herein, CNF/BG composites were prepared by sintering as-spun PAN/precursor composite nanofibers in nitrogen atmosphere at temperatures of 800, 1000 or 1200 °C for different times. The resulting CNF/BG composites were comprehensively and quantitatively analyzed with characterizations including TEM, SEM-EDXA and XRD etc., to expose the dependence of growth mechanism of BG nanoparticles on sintering temperature and time.
![]() | ||
Fig. 2 Morphologies of (a) the as-spun and (b) pre-oxidized PAN/BG precursor composite nanofibers observed by SEM and TEM (the inserts). The insets are the diameter distribution of nanofibers. |
Fig. 3 shows morphology evolution of CNF/BG-800 with various sintering time periods at 800 °C. Upon high-temperature sintering, PAN was carbonized into carbon and BG precursors would transform into BG. In this transition, the fiber diameter would decrease gradually with the loss of non-carbon elements. As shown in Fig. 3(a1), (b1) and (c1), the average diameters of CNF/BG-800 composites decreased with longer retention time at 800 °C, and were smaller than that of the pre-oxidized nanofibers. Some nano-scaled grains began to appear on the surfaces of CNFs, and the particle number increased slowly with sintering time prolonging. After 1 h sintering at 800 °C, numerous grains with average size of 28 nm were observed inside CNFs (Fig. 3(a2)), although the fiber surface remained smooth. The grains were identified BG nanoparticles as found in our previous work.31 When the sintering time reached 2 h, the emerged grains increased the surface roughness of CNF/BG-800 (Fig. 3(b1)), and the grains gradually grew into bigger size at the third hour of sintering (Fig. 3(c1)). Along with BG grains showing the tendency of migration towards outer surfaces of fibers, the TEM images revealed that those grains inside CNFs also grew bigger with broadened size distribution and declined distribution density. This change apparently met the quality conservation laws. It could be summarized from the results that 3 h sintering at 800 °C was not effective enough in propelling BG nanoparticles to the surface of CNFs.
When the sintering temperature was raised to 1000 °C, completely different morphologies were detected in resulting CNF/BG-1000 composites. As shown in Fig. 4, a mass of BG nanoparticles sprung up on surfaces of CNFs, with average sizes increasing from 45 to 75 nm as sintering time proceeding from 1 h to 3 h. The BG nanoparticles were spherical initially, and then evolved to irregular shaped aggregates of bigger sizes by fusion of individual nanoparticles. The particle sizes and distributions of BG nanoparticles inside CNFs changed accordingly. Bigger and denser BG grains were found on fiber surfaces, at the same time, bigger but less denser BG grains were found inside the fibers. These implied the nanoparticles got more thermodynamical driving force by raising sintering temperature to 1000 °C, which promoted them to migrate from the inside to the outside of the nanofibers. The driving force was associated with the decrease in chemical potential due to the reduction of interfacial area between BG nanoparticles and CNF matrix. After migration to the surface of CNFs, BG nanoparticles existed in a more stable thermodynamical state. With sintering time prolonging, the morphology of BG nanoparticles evolved from spherical to irregular shape due to the random and isotropic diffusion, migration, collision and mergence of small nanoparticles. It could be seen clearly in Fig. 4(b1), (b2), (c1) and (c2), some nanoparticles have merged together, but the contours of the original nanoparticles were still apparent as marked with blue rectangles. Some nanoparticles just collided together as marked with red ellipse, tending to form contact angle which built up the prerequisite for further mergence. The differences in morphology evolution of CNF/BG-1000 and CNF/BG-800 indicated that the formation mechanism of BG nanoparticles should be different. In addition, the degree of carbonization of PAN at 800 °C and 1000 °C was also different. In comparison with CNFs carbonized at 800 °C, the CNFs carbonized at 1000 °C were apparently thinner, indicating the loss of non-carbon elements to higher degrees.
By further increasing the sintering temperature to 1200 °C, the morphology of CNF/BG-1200 was found definitely different from the two previous cases. As seen from TEM images in Fig. 5, the majority of BG nanoparticles had migrated to the outside of nanofibers and only a few of particles stayed inside the nanofibers. The growth and migration of nanoparticles were significantly accelerated because the thermodynamical driving force went stronger at higher temperature. The BG nanoparticles of CNF/BG-1200 presented cuboid-like geometry at sintering time points of 1 h and 2 h (Fig. 5(a1) and (b1)), which were different from those spherical BG nanoparticles in cases of CNF/BG-800 and CNF/BG-1000, implying the different formation. mechanisms of BG nanoparticles at different sintering temperatures. Along with longer sintering time, significant fusion between nanoparticles took place to result in large size (107 nm), irregular aggregates (Fig. 5(c1)). Noticeably, the diameters of CNFs resulting from carbonization at 1200 °C were further decreased in comparison with those carbonized at 800 and 1000 °C. In addition to the increased thermodynamical driving force that accompanying with the increasing temperature, the continuously decreasing fiber diameter was believed another reason to cause BG nanoparticles migrating to fiber surface. During carbonization, the microstructure of PAN-based carbon fiber transformed from two-dimensional disordered graphite structure to three-dimensional laminar structure with loss of non-carbon elements, leading to structure densification at the same time.35,36 The densification was determined by both the rate of denitrogenation and the rearrangement of carbon atoms, which were proportional to sintering temperature. The denser carbon network might create stronger expelling force to “squeeze out” BG nanoparticles due to their chemical incompatibility.
![]() | ||
Fig. 6 HR-TEM images (left columns) and FFT patterns (right columns) of BG nanoparticles in CNF/BG-800 with various sintering time periods of (a1 and a2) 1 h, (b1 and b2) 2 h and (c1 and c2) 3 h. |
![]() | ||
Fig. 7 HR-TEM images (left columns) and FFT patterns (right columns) of BG nanoparticles in CNF/BG-1000 with various sintering time periods of (a1 and a2) 1 h, (b1 and b2) 2 h and (c1 and c2) 3 h. |
![]() | ||
Fig. 8 HR-TEM images (left columns) and FFT patterns (right columns) of BG nanoparticles in CNF/BG-1200 with various sintering time periods of (a1 and a2) 1 h, (b1 and b2) 2 h and (c1 and c2) 3 h. |
![]() | ||
Fig. 9 XRD patterns of (a) CNF/BG composites and (b) pure BGs sintered at different temperatures for 3 h. |
Accordingly, the well-defined lattice fringes with an average d-spacing of 0.71 nm found in the HR-TEM images of CNF/BG-1000 (Fig. 7), could confirm the formation of the (1 0 0) lattice plane of wollastonite structure.43,44 Similarly, the well-defined lattice fringes with an average d-spacing of 1.41 nm found in the HR-TEM images of CNF/BG-1200 (Fig. 8) could also confirm the formation of the (1 0 1) lattice plane of pseudowollastonite structure.45,46 Wollastonite and pseudowollastonite had different crystal structures. The former possessed chain structure while the latter possessed ring structure in the space of crystals, which resulted in the differences in the formation temperatures and the relating thermal stability. In addition, all these findings were consistent with corresponding FFT patterns. In Fig. 7, the crystallized lattices were observed disordered and the FFT patterns showed concentric diffraction rings associated with some diffuse diffraction spots, indicating wollastonite-type BG nanoparticles in CNF/BG-1000 mainly being in the polycrystalline form. From Fig. 8, highly ordered lattices with uniform interplanar spacing were observed and the corresponding FFT patterns displayed specific discrete diffraction spots, suggesting pseudowollastonite-type BG nanoparticles in CNF/BG-1200 being in the single crystalline form. In literatures, evidences showed that pseudowollastonite phase was a stable phase at high temperature, and the transformation from wollastonite to pseudowollastonite phase took place at temperatures over 1100 °C.47,48 Therefore, in a word, BG nanoparticles were in the form of polycrystalline wollastonite when the sintering temperature was 1000 °C, and transformed into more stable pseudowollastonite form at higher sintering temperature of 1200 °C.
In Ostwald ripening mechanism, a characteristic stated is the formation of spherical-type particles, for the material diffusion in the atomistic level was isotropic as mentioned in our previous study27 and literature.26,49,50 In coalescence mechanism, whereas, the feature is the formation of irregular-shaped particles as a result of anisotropic and random diffusion–collision of sub-particles.25,51,52 Therefore, the growth of BG nanoparticles with well-distributed spherical shape along CNFs was dominated by Ostwald ripening mechanism at the sintering temperature of 800 °C and the initial stage of the sintering temperature of 1000 °C. As the sintering going on, the spherical BG nanoparticles in CNF/BG-1000 began to merge into irregular-shaped particles, which meant coalescence had taken part in the morphology evolution. While in the case of CNF/BG-1200, the initial cuboid-like BG nanoparticles was apparently controlled by coalescence mechanism from the beginning at the sintering temperature of 1200 °C, which led to irregular-shaped particles in large size liable to fall off.
Theory analysis showed that each mechanism would result in obvious scale invariant distribution functions for the particle size, F(d/〈d〉) vs. d/〈d〉, in the asymptotic limit. In addition to using geometry evolution of nanoparticles, therefore, the size distributions of BG nanoparticles obtained under different conditions could be supplemental evidences to deduce the mechanism that engendered coarsening. Accordingly, the size distribution function of BG nanoparticles in CNFs was analyzed by normalizing the cluster size frequency distribution at each time point as shown in Fig. 10. It was found that all the data did not collapse onto a single curve, as expected for a self-similar coarsening process. As shown in Fig. 10(a) and (b), the experimental particle distribution of BG nanoparticles obtained at 800 °C and at earlier time period (≤2 h) of 1000 °C showed the best agreement with the diffusion-controlled Ostwald ripening equation (eqn (1)):53,54
![]() | (1) |
![]() | (2) |
The relative contributions of both Ostwald ripening and coalescence process could also be evaluated from the first and third moments of the nanoparticle size distribution, μ1 = d3/dh and μ3 = 〈d〉/d3, where 〈d〉 = ∑di/n is the arithmetic mean radius, d3 = (∑di3/n)1/3 is the cube-mean radius, Rh = n/(∑(1/Ri)) is the harmonic mean radius. Both moments were unity for monodisperse particles resulted from condensation, and the growth of the particles was only predominated by coalescence mechanism when μ1 > 1.25 and μ3 < 0.905.57,58 If 1 < μ1 < 1.25 and 0.905 < μ3 < 1, the growth of nanoparticles was controlled by both mechanisms. Fig. 11 shows the first (μ1) and third (μ3) moments of the size distribution function of BG nanoparticles in CNFs. For example, the moments of the particle size distribution at t = 3 h at 1000 °C were μ1 = 1.64 and μ3 = 0.81 (Fig. 11(b)), implying that the coalescence should play a dominating role in the growth of BG nanoparticles at this stage. Therefore, the calculations confirmed that the growth of BG nanoparticles at 800 and 1200 °C was predominately controlled by Ostwald ripening and coalescence mechanism, respectively, while the growth of BG nanoparticles at 1000 °C was firstly controlled by Ostwald ripening mechanism followed by coalescence mechanism with longer sintering time.
Based on all the above results, evolution mechanism of BG nanoparticles in CNFs sintered at various conditions was proposed as illustrated in Fig. 12. It was found that the sintering temperature and sintering time were the determinant factors in controlling the morphology and structure of BG nanoparticles in CNFs. The quantity of BG nanoparticles migrating onto the surfaces of CNFs was proportional to the sintering temperature and sintering time, which obviously dominated the size and the size distribution of BG nanoparticles onto CNFs. This was due to the enhancement of the thermodynamical driving force for particle migration by increasing sintering temperature or prolonging sintering time. At low sintering temperature (e.g. 800 °C), BG nanoparticles possessed small size and most of them remained inside CNFs. At high sintering temperature (e.g. 1200 °C), BG nanoparticles with big size migrated onto the surfaces of CNFs. Rationally, BG nanoparticles grew into bigger size with the increasing of sintering time. However, sintering temperature was identified the primary factor in controlling the crystallization behavior and growth mechanism of BG nanoparticles. At low sintering temperature (e.g. 800 °C), only numerous BG nanoparticles with amorphous structure were detected, and the growth of BG nanoparticles could be well explained by the ripening mechanism. At high sintering temperature (e.g. 1200 °C), BG nanoparticles emerged in the cuboid-like geometry, mainly composed of pseudowollastonite (Ca3(Si3O9)) single crystals, which were dominated by migration–coalescence mechanism. However, the growth of BG nanoparticles was controlled by both ripening and migration–coalescence mechanisms at the sintering temperature of 1000 °C. The resulting BG nanoparticles were composed of wollastonite (β-CaSiO3) polycrystals. In the process of BG nanoparticles formation and migration, the structure transformation and densification of CNF matrix along with carbonization were another possible factor playing important role in the morphology and structure evolution of BG particles. When more three-dimensional laminar structure in CNFs were formed along with the significant structure densification, more BG-related materials should be expelled under thermodynamical driving force generated at higher sintering temperature or longer sintering time.
This journal is © The Royal Society of Chemistry 2014 |