Direct ink writing core-shell [email protected] scaffolds with tailorable shell micropores favorable for optimizing physicochemical and biodegradation properties

Journal Pre-proof Direct Ink Writing Core-shell Wollastonite@Diopside Scaffolds with Tailorable Shell Micropores Favorable for Optimizing Physicochemical and Biodegradation Properties Jianhua Shen, Xianyan Yang, Ronghuan Wu, Miaoda Shen, Fengling Lu, Feng Zhang, Zhao Chen, Xiaoyi Chen, Sanzhong Xu, Changyou Gao, Zhongru Gou

PII:

S0955-2219(19)30670-3

DOI:

https://doi.org/10.1016/j.jeurceramsoc.2019.09.049

Reference:

JECS 12759

To appear in:

Journal of the European Ceramic Society

Received Date:

12 June 2019

Revised Date:

27 September 2019

Accepted Date:

29 September 2019

Please cite this article as: Shen J, Yang X, Wu R, Shen M, Lu F, Zhang F, Chen Z, Chen X, Xu S, Gao C, Gou Z, Direct Ink Writing Core-shell Wollastonite@Diopside Scaffolds with Tailorable Shell Micropores Favorable for Optimizing Physicochemical and Biodegradation Properties, Journal of the European Ceramic Society (2019), doi: https://doi.org/10.1016/j.jeurceramsoc.2019.09.049

This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier.

Direct Ink Writing Core-shell Wollastonite@Diopside Scaffolds with Tailorable Shell Micropores Favorable for Optimizing Physicochemical and Biodegradation Properties

Jianhua Shen1, Xianyan Yang1, Ronghuan Wu2, Miaoda Shen2, Fengling Lu1, Feng Zhang3, Zhao Chen4, Xiaoyi Chen4, Sanzhong Xu2*,Changyou Gao1, Zhongru Gou1* 1

Bio-nanomaterials and Regenerative Medicine Research Division, Zhejiang-California International Nanosystem

Institute, Zhejiang University, Hangzhou 310058, China. 2

Department of Orthopedics, The First Affiliated Hospital, School of Medicine of Zhejiang University, Hangzhou

310003, China. Department of Stomatology, Children’s hospital, School of Medicine of Zhejiang University, Hangzhou 310003,

3

ro of

China. Department of Orthopedics, Zhejiang Provincial People’s Hospital, People’s Hospital of Hangzhou Medical

4

-p

College, Hangzhou 310014, China.

*Corresponding author

re

Zhongru Gou, Ph.D.

Zhejiang-California International Nanosystems Institute

Tel: (+86) 571-8820 8353 Fax: (+86) 571-8697 1539

na

E-mail: [email protected]

lP

Zhejiang University, Yuhangtang Road 866#, Hangzhou 310058, China

Sanzhong Xu, M.D., Prof.

ur

Department of Orthopedics, the First Affiliated Hospital School of Medicine of Zhejiang University

Jo

Qunchun Road 79#, Hangzhou 310009, China E-mail: [email protected]

Abstract Additive manufacture has recently been proposed as a versatile process for fabricating porous bioceramic scaffolds 1

for bone repair and tissue engineering; however, to control or tailor the biodegradation of the porous biomaterials is still a challenge. Here, the core-shell-structured biphasic bioceramic porous scaffolds with tailorable ion release and biodegradation were prepared by direct ink writing technique with coaxially aligned bi-nozzle system. Our method employed rapidly gelling filaments of wollastonite (CSi) and diopside without and with Zn or Sr doping (Dio, ZnDio, SrDio) derived from bi-flow of sodium alginate-loaded bioceramic slurries, and varying the powder slurry design made it easy to create core-shell struts (e.g. CSi@Dio, CSi@ZnDio, CSi@SrDio) with adjustable bioceramicphase distribution. It was found that the Zn- or Sr-doping could readily adjust the mechanical strength and biodegradation rate in the early stage. Furthermore, when 30% organic microspheres were pre-mixed into the powder slurry, the controllable high-density micropores could be introduced into the shell layer after sintering (e.g.

ro of

CSi@Dio-p, CSi@ZnDio-p, CSi@SrDio-p), and thus permeability is maximally tuned and favorable for ion release through the porous shell layer. This new core-shell direct ink writing strategy can be used to fabricate a variety of biphasic bioceramic scaffolds with adjustable physicochemical properties which could be potentially beneficial for improving biological performance and bone repair in situ.

-p

Keywords: Core-shell strut; Biodegradation; Mechanical strength; Directing ink writing; bioceramic scaffolds.

re

1. Introduction

It is known that there are a great amount of bone fractures or trauma suffering incomplete healing due to bone loss,

lP

failed fixation, infection, and inadequate vascularization[1-3]. Although bone tissue has the ability to self-renewal soon after injured, the fibrous connective tissue may grow into the bone defect once the defect is beyond the critical size or accompanied with chronic infection, hindering the bone healing process. Therefore, the bioactive scaffolds

na

with appreciable porosity and pore interconnectivity become the main replacement for grafting in bone regenerative medicine[4,5]. Basically, tissue ingrowth through three-dimensional(3D) interconnected pores provides enhanced mechanical interlocking between host tissue and porous biomaterials[6]. Also, pore interconnectivity is beneficial for

ur

nutrient transport to and from the core of the implants, which is very crucial for proper vascularization[7]. Hence, the scaffold design to support bone formation largely determines the success of bone repair.

Jo

Currently, the porous bioceramics with controllable biodegradation rate and appreciable mechanical strength are

thought to be the promising candidates for bone regenerative medicine[8,9]. Such biomaterials may release biologically active ions in body fluid and react with the host at surface through the surface microstructure to induce and promote the new bone formation[10]. Subsequently, the porous bioceramics will degrade gradually in the process of tissue regeneration, and eventually be completely replaced by newly formed bone tissue. It is expected that, in particular, the rate of material biodegradation need to match the rate of bone tissue formation so that a room for matrix deposition and new bone ingrowth can be provided[11]. 2

Over the past two decades, special and increasing attention has been paid to the silicate-based biomaterials as potential biomaterial for bone repair due to their appreciable bioactivity. Most modern research into the use of silicate-based biomaterials for bone repair dates back to the pioneer work by Hench in the 1970s, who first demonstrated the bioactivity and hydroxyapatite re-mineralization nature of CaO-SiO2-based bioactive glasses[12]. Ca-silicate bioceramics such as wollastonite (CaSiO3; CSi) and diopside (CaMgSi2O6; Dio) are the hot point of researches because of their good biocompatibility[13-15]. Meanwhile, these bioceramics showed good osteoconduction and/or osteoproduction due to their remarkable bioactivity and different bioresorbability in vivo[16,17]. However, it seems that any single-phase biomaterial among these silicate bioceramics is not able to be an ideal candidate for stimulating in situ bone regeneration and repair due to their mismatching biodegradation

ro of

and/or mechanical evolution requirement with time[5]. For instance, the porous CSi scaffolds readily undergo pore collapsing through rapid biodegradation in vivo[18]. In contrast, it is indicated that the Dio scaffolds are significantly less soluble than the other Ca-silicate bioceramics, and displays slower biodegradation rate in vivo in comparison with the new bone tissue ingrowth[17,19]. Sainz et al. reported that the CSi/Dio composites displayed a relatively

-p

high bioactivity[20]. Our recent studies showed the dilute Mg-substituted CSi could readily reinforce the Dio

scaffolds, and the secondary phase in the porous scaffolds can accelerate the biodegradation but retard mechanical

re

decay[21-23]. Unfortunately, such mechanically mixing hybrid cannot easily facilitate to tailor the physicochemical and biological properties of the bioceramic composite. One way to control the biodegradation rate and bioactivity of

lP

such composite scaffolds is reached by increasing one component while decrease the other component in the composites to controlling the biodegradation kinetics of the whole biomaterials. However, it is a great challenge for

na

simultaneously tailoring the diverse properties of the porous composites. On the other hand, there are many advanced methods to fabricate the porous scaffolds, such as additive manufacture, microsphere- or ice-template process and so on[24-29]. It is demonstrated that the closely packed

ur

microsphere- or ice-template approach can produce highly ordered or hierarchical porous structures with outstanding mechanical properties. Despite the many advances in fabricating bioceramic scaffolds by these methods, additive

Jo

manufacture remains an attractive and popular choice of bioceramic scaffolds for bone regeneration[30]. The many qualities of additive manufacture techniques include its periodic pore structure, limited pore-strut defect, flexible scaffold morphology, and adjustable geometrical parameters. Interestingly, most research involving the use of additive manufacture for porous bioceramics is the direct ink writing (DIW) technique, which may help to fabricate the ideal porous scaffolds by extruding inorganic powder slurry through the printing nozzle and accurately control the pore structure, pore size and porosity through computer[31]. Despite continuing advances of DIW technique, however, there is a challenge for the bioceramic scaffolds with tailorable physicochemical properties with time. 3

Our previous work have designed the core-shell-structured bioceramic granules and examined the component distribution on the material properties and bone formation in vivo[32-34]. In the initial study, the different bioceramic components could be readily integrated into the core or shell layer of biphasic granules with tunable core or shell composition through the coaxially aligned capillary system[32]. Other studies examined the different component microstructures on the bio-dissolution and ion release behavior in vitro and osteogenic capability in vivo in diverse bone defect models. Intriguingly, it was indicated that the foreign ion doping or micropore structures in core or shell layer could tune the bone repair efficacy[33,34]. The results from these studies emphasized the fact that changes in the component distribution and microstructure of specific component have significant effects on the final biological performances of the bioceramics.

ro of

Based on these concerns, we try to design the CSi‒Dio biphasic bioceramic scaffolds with core-shell-structured pore struts. In order to better understand the relationship between chemical composition and microstructure in pore struts and expected mechanical and biodegradable properties in vitro, the foreign ion doping and/or high-density micropore design were studied by doping zinc or strontium into Dio powder and adding organic microsphere

-p

porogen in the Dio bioceramic slurries. This study showed that the tailorable component distribution and porous

and biodegradable properties of bioceramic scaffolds.

2.1 Chemicals and Materials

lP

2. Materials and methods

re

microstructures in pore struts of biphasic bioceramic scaffolds is beneficial for simultaneously tuning the mechanical

The reagent-grade inorganic salts, trishydroxymethylaminomethane (Tris) and tetraethyl orthosilicate (TEOS)

na

were bought from Sinopharm Reagent Co., Shanghai, and used without further purification. Polyvinyl alcohol (PVA) was purchased from Sigma-Aldrich. Tris was used to prepare to the 0.05 M Tris buffer (pH ~7.25).

ur

2.2 Preparation of CSi and Dio bioceramic powders The CSi powder was synthesized by a wet-chemical co-precipitation method as described previously[21]. The as-

Jo

calcined CSi powders were CSi powders were ground by zirconia ball media in ethanol for 4 h. The particle size of the resulting powders was below 5 μm. The Dio powders doped with 4% Sr or Zn (4 mol% Mg was substituted by Zn or 4 mol% Ca was substituted by Sr; ZnDio or SrDio) were synthesized by a conventional sol-gel method[35]. Firstly, Ca(NO3)2·4H2O, Mg(NO3)2·6H2O, and Zn(NO3)2·6H2O or Sr(NO3)2 in stoichiometric ratio were dissolved completely in absolute ethanol and then TEOS was added. The solution was stirred rapidly for 2 h. The pH value

4

(10.5~11.0) was maintained by adding dropwise with NH 3∙H2O. Then, the sol was stirred for 4 h and aged at 80℃ for 24 h. After that, the gel was dried at 120oC overnight and followed by calcining at 1100℃ for 2 h. Moreover, the pure Dio powder was synthesized in the absence of Sr2+ and Zn2+ ions in the sol reactant while the other conditions remained the same. The Dio, SrDio and ZnDio powders were ground using zirconia ball media in ethanol for 6 h. The particle size of the resulting powders was below 5 μm. 2.3 Preparation of bioceramic scaffolds The bioceramic powder slurries for layer-by-layer writing of the porous scaffolds with core-shell struts were

ro of

prepared by mixing 4.5 g bioceramic powders and 4.0 g of 10 wt% PVA solution. As for the scaffolds with CSi core and Dio shell in pore struts (denoted as CSi@Dio), the CSi and Dio slurries reached the bi-layer nozzles through

different microtubes, core–shell pastes of bioceramic slurries were formed (see Fig.1 B). The interior diameters of core and shell nozzles were 400 μm and 900 μm, respectively. And the spacing between filaments was designed as

-p

500 μm. Meanwhile, the layer thickness between the two adjacent layers was 1000 μm. Then the CSi@Dio paste for

re

layer-by-layer writing as porous scaffolds (6 × 6 × 8 mm) was prepared by the DIW equipment. The scaffolds were manufactured by modifying the deposition angle from 0° to 90° after two printed layers as shown in Scheme 1. The

lP

samples were dried at 80oC for 24 h and then sintered at 1250oC in air atmosphere using the heating schemes (the heating rate is 2oC/min while maintaining at 320oC and 500oC for 45 and 60 min, respectively) and held at the target

na

temperature for 3 h, followed by cooling naturally.

In order to investigate the effect of shell-layer component on the physicochemical properties of the scaffolds, the

ur

scaffolds including CSi@ZnDio and CSi@SrDio were also prepared while the other conditions remained the same. Similarly, the CSi@CSi, Dio@Dio scaffolds were prepared as the procedures described above.

Jo

Moreover, the ~5 μm polystyrene (PS) microspheres were mixed with the Dio, SrDio or ZnDio-based slurry with a

PS-to-bioceramic powder mass fraction of 30%. As the slurries reached the coaxially aligned nozzles through different microtubes, the paste for layer-by-layer writing as porous scaffolds of CSi@Dio-p, CSi@SrDio-p, CSi@ZnDio-p (6 × 6 × 8 mm) was prepared, while the other conditions remained the same. 2.4 Microstructure and phase analyses The chemical composition of the bioceramic powders was measured by inductively coupled plasma-optical 5

emission spectrometry (ICP-OES; Thermo). The phase-composition characterization was performed by X-ray diffractometer (XRD; Rigaku D/max-rA) with a monochromatic CuKα RADIATION (λ=1.5406Å) operating at 30 kV/20 mA. The XRD patterns were recorded in steps of 0.02° intervals with a counting time of 0.5 s at each step. Scanning electron microscopy (SEM; JEM-6700F) with energy-dispersive spectroscopy (EDX) was used to characterize the Au-coated (fracture) surface of the porous scaffolds at an accelerating voltage of 10 kV. 2.5 Mechanical testing Porous scaffolds (n=3; 6 × 6× 8 mm) were prepared for compressive strength. Mechanical testing was performed

ro of

using a static mechanical test machine (Instron 5566) and a 10 kN load cell, and samples were tested at a constant cross head displacement rate of 0.5 mm/min. 2.6 Degradation evaluation in vitro

For evaluation of degradation, the porous scaffolds (n=12; 6 × 6 × 8 mm) were immersed in Tris buffer with an

-p

initial pH 7.4 and S/L ratio of 1.0 g/50 ml for 0-8 weeks. After immersing for 4 h, 1, 3, 5, 7 and 14 d, the supernatant

re

(0.5 ml) was extracted for ICP-OES measurement and the fresh buffer (0.5 ml) was added to keep solution volume constant. After immersion for 2, 4, and 8 weeks, the scaffolds (n=3) were rinsed with ethanol and then dried at 100oC

lP

for 12 h, and weighed (Wt). The weight loss at time t was expressed as the equation: Weight decrease=Wt /W0×100%. The compressive strength of the immersed scaffolds was determined by Instron testing machine.

na

2.7 Biomimetic apatite re-mineralization in SBF in vitro

The CSi@Dio scaffolds (~0.20 g; 6 × 6 × 8 mm; n = 3) were immersed in 28.8 ml SBF at 37oC and monitored the

ur

formation of carbonated apatite on the surface of samples[36]. After soaking for 7 d, the scaffolds were washed with ethanol and observed using SEM and a local chemical analysis was carried out by EDX. Prior to examination, the

Jo

samples were coated with a thin layer of gold. 2.8 Statistical analysis

All the data were expressed as mean ± standard deviation and analyzed with the one-way ANOVA. In all cases the

results were considered statistically significant with a p-value less than 0.05.

3. Results 3.1 Primary characterization of the bioceramic powders 6

Table 1 reports the ICP data for the ZnDio and SrDio powders. It is showed that the Zn-substituting-Mg ratio and Sr-substituting-Ca ratio was 3.31% and 4.78%, respectively. These measured data were similar to the theoretical substituting ratio of 4%, which may be attributed to the approximate ion radius between Ca2+ (100 pm) and Sr2+ (118 pm) and between Mg2+ (72 pm) and Zn2+ (74 pm), respectively. Figure 1A shows the XRD pattern of CSi and Dio powders with and without foreign ion doping. It can be found that these CSi and Dio powders maintained respectively the wollastonite phase (PDF# 42-0547) or diopside phase (PDF#86-0932). The peak in powder XRD patterns recorded at 29.7o/2θ was subjected to magnification to determine

ro of

the influence of Sr or Zn substitution in Dio (Fig. 1B). Indeed, a slight peak shift for this (-2, 2, 1) crystal plane to higher or lower value was observed for the SrDio and ZnDio powders, respectively. It implies the Dio was exactly substituted by the foreign ions to a certain extent.

3.2 Microstructure and mechanical characterization of the bioceramic scaffolds

-p

Figure 2 reports the SEM images (low-magnification) of the fracture surface for the porous scaffolds after

re

sintering. SEM analysis from the internal side-wall observation revealed that the macroporous architecture was maintained in all groups of scaffolds, and especially the scaffolds with Dio shell layer (i.e., Dio@Dio, CSi@Dio,

lP

CSi@Dio-p) showed more appreciable macropore dimension (Fig. 2B‒D). Comparing the single-phase (i.e., CSi@CSi, Dio@Dio) with the biphasic (i.e., CSi@Dio, CSi@Dio-p) bioceramic scaffolds, it is possible to appreciate

na

how obviously the fracture microstructure of the pore struts changed in the later due to the presence of definite coreshell boundaries. In fact, after intruding different bioceramic slurries through the core-shell nozzles, a bright core-

ur

shell-structured strut was observed in the sintered bioceramic scaffolds (Fig. 2C, D). Pore-strut SEM observation and corresponding EDX analysis of biphasic bioceramic scaffolds (Fig. 3 & Tab. 2)

Jo

depict the different composition and core-shell microstructures with distinct boundaries; however, lack of uniformity in the micropores is inevitable in the shell layer. It is revealed that the CSi core was readily coated by the external shell of Dio, ZnDio, or SrDio (Fig. 3A, C, E). Furthermore, the microstructure of shell layer in the CSi@Dio, CSi@ZnDio, and CSi@SrDio scaffolds displayed subtle differences with that of CSi@Dio-p, CSi@ZnDio-p, and CSi@SrDio-p scaffolds, in which the high-density micropores were obvious and moreover their distributions were attributed with the embedded PS microspheres in the shell-layer powder slurries in advance (Fig. 3B, D, F). This is 7

also confirmed by the magnified SEM images as shown in Figure 3A1 and 3B1, in which some significant differences were observed at the fracture surface of the pore struts in CSi@Dio and CSi@Dio-p samples. More spherical micropores (white arrows) were observed in the CSi@Dio-p scaffolds. Also, EDX spectra revealed the presence of Ca, Si, and O atoms, and in particular the Dio shell layers may be differentiated from the CSi core layer based on the peek strength for Mg atoms (see Fig. 3G, H(insets)). Such EDX analysis relatively ensured presence of Mg atoms in the Dio shell layer. The compressive test for the porous scaffolds with different composition and microstructures in shell layer was

ro of

determined through universal testing mode and the corresponding results are presented in Figure 4. As shown in Figure 4A, the pure CSi@CSi and Dio@Dio scaffolds showed significantly different compression resistance (14.2 MPa vs 26.5 MPa) after undergoing same sintering treatment. As it was expected, the presence of high-density micropores in the Dio, SrDio or SrDio shell layer had a strong influence on the strength of the scaffolds. The

-p

compressive strength of the scaffolds without shell-layer micropores (i.e. CSi@Dio, CSi@ZnDio, CSi@SrDio) were

re

nearly 2-fold higher than that of the scaffolds with shell-layer micropores (i.e. CSi@Dio-p, CSi@ZnDio-p, CSi@SrDio-p). On the other hand, it is worth mentioning that, as for the biphasic bioceramic scaffolds, the Sr doping

lP

in Dio shell layer may be helpful for enhancing the compression resistance in comparison with the Zn doping. Moreover, the Young’s modulus of scaffolds without and with shell-layer micropores were respectively determined

na

as over 250 MPa and 220 MPa (Fig. 4B), which ensured relatively appreciable values close to the cancellous bone. The stress vs. strain curves recorded from all of scaffold compositions determined from the compressive strength

ur

measurement were showed in Figure 4C‒F. The stress-strain curves demonstrated a steady upsurge but a rapid drop (i.e. elastic deformation) for the bioceramics scaffolds without shell-layer micropores; nevertheless, a plastic

Jo

deformation for the scaffolds with shell-layer micropores was noticed as a function of strain. 3.3 Bio-dissolution and re-mineralization evaluation in vitro To further investigate the potential biodegradation behavior and strength evolution of the scaffolds in vivo, the

sintered scaffolds were immersed in Tris buffer from 2 to 8 weeks, to simulate the pH microenvironment in body. From the mass decrease graphics shown in Figure 5A, the lowest and highest curves observed were ascribed to the pure CSi@CSi and Dio@Dio bioceramic scaffolds, which displayed the fastest and slowest bio-dissolution, 8

respectively, in comparison with the other biphasic bioceramic scaffolds in Tris buffer. The Dio@Dio scaffolds showed only ~1.8% mass loss after 8 weeks, but that for the CSi@CSi scaffolds (~9.6%) was nearly 5-fold higher than the Dio@Dio. On the other hand, the shell-layer micropores in the biphasic CSi@Dio-p scaffolds contributed to a faster bio-dissolution than the corresponding CSi@Dio scaffolds. However, it is interesting that, the Sr and Zn doping into the Dio component would accelerate and inhibit the bio-dissolution of the biphasic bioceramic scaffolds, respectively. After immersion for 8 weeks, the mass loss CSi@SrDio-p and CSi@ZnDio-p was 7.92% and 4.25%, while the mass loss of the CSi@Dio-p was 5.98% at the end of the immersion period.

ro of

The strength evolution of the porous scaffolds during immersion in Tris buffer was determined by mechanical test, as shown in Figure 5B. The CSi@CSi scaffolds showed the lowest initial strength before immersion, and displayed a rapid decrease by nearly 46% within 2 weeks, from ~14.2 MPa to ~7.7 MPa, and then maintained a very slow

reduction up to 8 weeks (~7.4 MPa). In contrast, the Dio@DiO scaffolds changed very slowly with limited strength

-p

decay from ~26.5 MPa to 24.5 MPa during the whole immersion stage. As for the three groups of biphasic

re

bioceramic scaffolds without and with shell-layer microspores, the compressive strength were decreased mildly with time, and the scaffolds without micropores maintained higher strength than the counterparts, which tended to exhibit

lP

good consistency with their initial strength.

ICP analysis of the aqueous medium (Fig. 6) depicted the ion concentration profiles during immersing the

na

bioceramic scaffolds; however, lack of uniformity in the changes in ion concentration was inevitable. It is interesting that the highly stable Dio@Dio scaffolds showed very slow Ca and Si release but appreciable Mg release (Fig. 6A-

ur

C). Indeed, the highly biodegradable CSi@CSi displayed fast Ca and Si release, although the shell-layer micropores in the CSi@SrDio-p scaffolds readily contributed to more appreciable CSi core dissolution and ion release.

Jo

Evidently, such porous microstructures in the pore struts were always favorable for Ca and Si release in the three groups of scaffolds including CSi@Dio-p, CSi@ZnDio-p and CSi@SrDio-p, in comparison with their counterparts without shell-layer micropore structures (i.e. CSi@Dio, CSi@ZnDio, CSi@SrDio). Moreover, the higher specific surface area in the ZnDio or SrDio shell layer would contribute to the Zn or Sr release (Fig. 6E, F), although the Sr concentration was ten-fold higher than the Zn concentration in the similar immersion condition, possibly due to the significantly different bio-dissolution properties of the ZnDio and SrDio bioceramics. 9

SEM observation of the as-immersed samples (Fig. 7) showed the surface and fracture morphology of the pore struts in the CSi@Dio and CSi@Dio-p scaffolds with distinct microstructure evolution. Alongside, the obvious presence of micropore distribution and enlargement with time are also visible in the surface of CSi@Dio-p scaffolds. The high-magnification SEM images revealed the bio-dissolution-derived loose microstructure in the sparingly dissolvable Dio surface layer after 8 weeks (Fig. 7A-C). In particular, the diameter of the spherical micropores in the Dio layer became larger with the prolongation of immersion time (Fig. 7D-F). As for the fracture microstructures of the scaffolds (Fig. 7G-L), the core-shell interface of CSi@Dio biphasic bioceramic was discernible and the Dio shell

ro of

layer appeared to be more porous in the CSi@Dio-p scaffolds; however, the microstructure of the CSi core and Dio shell layer, on the other hand, was almost unaltered during the whole immersion stage.

To further determine the suitability of the bioceramic scaffolds for the application in bone repair, their surface

reactivity was tested by immersing SBF for 7 d (Fig. 8). It was observed that a dense precipitate layer was coated on

-p

the surface of the pore struts. Based on the quantitative EDX analysis, the Ca/P ratio for the surface layer was ~1.01-

re

3.84. It implies that a new Ca-phosphate re-mineralization layer was deposited onto the pore struts. Indeed, the highly stable ZnDio and highly biodegradable SrDio surface layer is more unfavorable for the rapid re-mineralization

lP

of apatite in such simulated biological aqueous medium.

4. Discussion

na

The controlled implantation of porous bioceramics in clinical practice started in 1980s with the use of Caphosphate bioceramics for bone filling[37,38]. Till now, the use of such porous implants as scaffolds for bone tissue augmentation and repair is expanding in orthopedic and dental areas[39]. This increase is in large part driven by the

ur

progress in bioceramic science, especially the new manufacturing technologies such as additive manufacture developed in the last two decades. However, much work is still required for bioactive ceramics to reach their full

Jo

potential because their bioactive, biodegradable, and mechanical properties, in particular the spatiotemporal evolution of their porous architecture with time are all of the major concerns[40-42]. Moreover, it is still difficult to tune the bioceramic biodegradation and bioactive ion release sufficiently favorable for stimulating new bone growth after implantation. In this study, we demonstrate how a facile and flexible one-step process can produce core-shell-structuring struts in biphasic bioceramic scaffolds with tunable shell composition and microstructures. We choose CSi and Dio to demonstrate the effectiveness of this method in fabricating component distribution-tuned porous bioceramics, as 10

these biomaterials exhibit good bioactivity, but show either faster or slower biodegradation rate than the new bone regeneration rate in vivo. Our experimental results confirmed that such CSi@Dio and CSi@Dio-p scaffolds of are readily endowed with time-dependent biodegradation and tuned bioactivity and strength evolution with time. This new strategy may be of benefit for preventing incomplete or retarding bone repair in pure CSi and Dio scaffolds, respectively; in particular it may reduce the risk of uncontrollable long-term inflammatory reaction triggered by the sparingly dissolvable microparticles from the conventional biphasic hybrid bioceramics. Previously, Paredes et al. have developed the hydroxyapatite scaffolds with enhanced strength and toughness by infiltrating with polycaprolactone into the hollow tubes (i.e. hollow strut) of the scaffolds [43]. Wu and collaborators have developed bioceramic scaffolds with novel strut structures including closely packed hotdog-like strut or hollow struts via direct

ro of

ink writing technique, and these porous architectures have been demonstrated some advantages favorable for vascularization and new bone tissue ingrowth [44-46]. However, these scaffolds are still difficult to tune the

biodegradation and bioactive ion release for stimulating new bone regeneration in the early stage after implantation. It has been reported that the Dio bioceramic is very biologically stable, comparable and even superior to the pure

-p

hydroxyapatite bioceramics[47]. However, Dio bioceramics exhibit good sintering properties and appreciable

mechanical strength in comparison with the other Ca-Mg-silicate bioceramics[17,19,48]. In this study, we found that

re

the mechanical and physicochemical properties of the Dio could be adjusted by Sr- or Zn-doping process. When 4% Ca or Mg was substituted by Sr or Zn in Dio bioeramic component, its phase purity was not influenced by the foreign

lP

ions, but the compressive resistance and biodegradation rate of the CSi@Dio scaffolds could be enhanced or reduced, respectively. This suggests the Dio shell layer in the scaffolds could be readily tuned by foreign ion doping.

na

On the other hand, once the organic microspheres were pre-mixed into the Dio bioceramic slurry, the increased microporous structures could be introduced into the Dio shell layer after sintering, and thus its permeability is tailored for rapid core-layer inorganic ion diffusion through the mcriopores in the shell layer. Thus, this core-shell

ur

printing design is beneficial for fabricating the porous bioceramic composites in which the biodegradation rate and bioactive ion release are readily predicted and tailored over time.

Jo

Sintering, as an important treatment stage in bioceramics processing, has also undergone significant modification

and some novel sintering technologies have been widely introduced. For example, hot pressure or isostatic pressure sintering is reported as successful approaches to control bioceramic grain growth during extremely high rates of heating process [49,50]. However, the sintering of bioceramic scaffolds cannot be underwent an additional pressure because the macroporous and microporous architecture is needed to be utmost extremely maintained in the scaffolds. Therefore, the mechanical improvement of bioceramic scaffolds is mainly relied on the sintering properties of 11

bioceramic compound itself and/or other additive agents. It is can be found from the SEM images that, as for the CSi@CSi and Dio@Dio scaffolds, the SEM observation only depicts the continuous distribution of grains and the core-shell interface is no more visible (Fig. 2A, B), In contrast, the distinct core-shell boundaries in the pore struts of the CSi@Dio and CSi@Dio-p scaffolds are easily discernible (Fig. 2C, D). This is mainly attributed to the significantly different sintering properties (or sintering shrinkage) of CSi and Dio phases. It is obvious that, as expected, the high-density micropores in the Dio shell layer without and with Zn or Sr doping can be observed clearly. This means the spherical micropores are retained well in the shell layer once the organic porogen is

ro of

volatilized during sintering process. Totally, these microstructure observations suggest that the initial microstructure of the pore struts in such core-shell bioceramic scaffolds is mainly relied on the addition of porogen and sintering properties of the bioceramic components.

-p

Our results demonstrate that the compressive resistance of the biphasic bioceramic scaffolds is related with the component distribution in the core and shell layer, foreign ion doping and porous microstructure in the pore struts.

re

Comparing the scaffolds with and without foreign ion doping in Dio shell layer, it is possible to appreciate how strongly the Sr doping in the Dio shell layer contribute to a significant enhancement of compressive strength and

lP

Young’s modulus (Fig. 4A, B). On the other hand, the strength evolution of bioceramic scaffolds immersed in Tris buffer is to illustrate how fundamental the choice of bioceramic composition and distribution is and how much the

na

micropore structure in shell layer can affect the strength. Coherently, in terms of structural properties, the strength and mass were almost changed consistently for the bioceramic scaffolds in Tris buffer, as shown in Figure 5.

ur

Although all bioceramic scaffolds underwent a decrease in mass, the stronger diminutions were observed on the scaffolds with spherical micropores in the shell layer. After immersion in Tris buffer for 8 weeks, SEM analyses

Jo

revealed that the bright increase of spherical micropores in the Dio shell layer of pore struts was observed (approx. 1.5-fold in diameter; Fig. 7D, F). This is also confirmed by ICP data reported in Figure 6, which summarizes some significant bio-dissolution and ion release from both the core and shell layer during immersion in such aqueous solution. Although in Tris buffer the CSi@ZnDio-p and CSi@SrDio-p scaffolds underwent respectively a stronger Zn release and a weaker Sr release in comparison with the CSi@ZnDio and CSi@SrDio scaffolds, their Ca release were higher in Tris buffer. This could be explained by the different bio-dissolution of CSi core and Dio shell layers. 12

Moreover, the surface reactivity in simulated biological medium is also evaluated to further determine the suitability of the bioceramic scaffolds favorable for bone regeneration and repair. It is known that the bioactivity of Ca-silicate and Ca-Mg-silicate bioactive materials is attributed to the formation of a bone mineral-like apatite layer on their surface in the (simulated) biological environments, thus a strong bond can form with the bone tissue[15,51]. Since the formation in vivo of an apatite film at the interface between the bone and the implant is necessary for the implant integration, the in vitro development of such a bone mineral-like film is usually considered as a fundamental preliminary requirement[52]. Previous studies have showed that the periods within 3–7 days for inducing apatite in

ro of

SBF is considered a necessary condition to convey bioactivity in vivo of bioactive materials, and thus the in vivo bioactivity can be precisely reproduced by apatite-forming ability in SBF [37,53]. It has been demonstrated that the pure Dio bioceramic exhibit appropriate bioactivity in vitro and in vivo [19]. Meanwhile, it is well agreed that the inorganic ion release is shown to convert faster and more completely to apatite in vitro and enhances new bone

-p

formation in vivo [14,16]. Based on the SEM/EDX analyses, the surface of the pore struts appear to be covered with

re

large and tightly packed Ca-phosphate granules (Fig. 8). This could be explained by the different microstructure and composition of the new precipitated phase observed on the bioceramic scaffolds. In fact, such microstructure is also

lP

evidently different with the bioceramic scaffolds after immersion in Tris buffer (Fig. 7A-F). When in Tris buffer, the surface compositions were mainly made of the highly crystalline CSi or Dio grains, and they appeared as low

na

cohesive agglomerates due to inappropriate sintering properties; in SBF, however, they were made of newly precipitated Ca-phosphate granules, and they were tightly bonded. That is, pH-buffered medium, for example, would

ur

be preferable for when the long-term signs of the evolution of the core and shell composition was studied. On the other hand, biomimetic mineral medium, would be more appropriate for investigating bioactivity phenomena, which

Jo

would normally take place on a shorter time scale. To conclude, the immersion solution seems to play a fundamental role too, affecting bio-dissolution or precipitation in the Dio shell layer of the biphasic bioceramic scaffolds.

4. Conclusions

In summary, our studies demonstrate that the coaxial bi-nozzles are versatile for core-shell CSi@Dio bioceramic scaffolds with selective Ca-silicate and Ca-Mg-silicate distribution in core or shell layer of pore struts via direct ceramic ink writing technique. In particular, the selectively foreign ion doping or high-density micropores in the Dio 13

component could be endowed by pre-doping in Dio powder or adding organic microsphere porogen in Dio slurry. These pretreatment steps are both favorable for tailoring the dynamic mechanical strength and bioactive ion release in the porous bioceramic composites with core-shell struts. Therefore, it is reasonable to assume that this facile and versatile strategy is promising for developing a wide range of biphasic bioceramic composite scaffolds with precisely tuned phase combination and distribution, tailorable microstructures, and controlled physicochemical properties for bone regenerative medicine applications. Acknowledgments

ro of

This work was supported by the National Key R&D Program of China (2017YFE0117700), the Science and Technology Department of Zhejiang Province Foundation (2017C37134, LGF18E020001), Zhejiang Provincial Natural Science Foundation of China (LY17H180010), and National Natural Science Foundation of China

Jo

ur

na

lP

re

-p

(81772311, 81871775).

14

References 1.

Loi F, Cordova LA, Pajarinen J, Lin T, Yao Z, Goodman SB. Inflammation, fracture and bone repair. Bone. 2016; 86: 119-30.

2.

Roberts TT, Rosenbaum AJ. Bone grafts, bone substitutes and orthobiologics: the bridge between basic science and clinical advancements in fracture healing. Organogenesis. 2012; 8: 114-24.

3.

Winkler T, Sass FA, Duda GN, Schmidt-Bleek K. A review of biomaterials in bone defect healing, remaining shortcomings and future opportunities for bone tissue engineering: The unsolved challenge. Bone Joint Res. 2018; 7: 232-43.

4.

Brydone AS, Meek D, Maclaine S. Bone grafting, orthopaedic biomaterials, and the clinical need for bone engineering. Proc Inst Mech Eng H. 2010; 224: 1329-43. Hutmacher DW. Scaffolds in tissue engineering bone and cartilage. Biomaterials. 2000; 21: 2529-43.

6.

Fielding G, Bose S. SiO2 and ZnO dopants in three-dimensionally printed tricalcium phosphate bone tissue

ro of

5.

engineering scaffolds enhance osteogenesis and angiogenesis in vivo. Acta Biomater. 2013; 9: 9137-48. 7.

Somo SI, Akar B, Bayrak ES, Larson JC, Appel AA, Mehdizadeh H, Cinar A, Brey EM. Pore interconnectivity influences growth factor-mediated vascularization in sphere-templated hydrogels. Tissue Eng Part C Methods. 2015; 21: 773-85.

Nommeots-Nomm A, Lee PD, Jones JR. Direct ink writing of highly bioactive glasses. J Euro Ceram Soc.

-p

8.

2018; 38: 837-844. 9.

Bose S, Roy M, Bandyopadhyay A. Recent advances in bone tissue engineering scaffolds. Trends Biotechnol.

re

2012; 30: 546-54.

10. Mourino V, Cattalini JP, Boccaccini AR. Metallic ions as therapeutic agents in tissue engineering scaffolds: an

lP

overview of their biological applications and strategies for new developments. J R Soc Interface. 2012; 9: 40119.

11. Kneser U, Schaefer DJ, Polykandriotis E, Horch RE. Tissue engineering of bone: the reconstructive surgeon's point of view. J Cell Mol Med. 2006; 10: 7-19.

Hench LL. The story of Bioglass®. J Mater Sci Mater Res. 2006; 17: 967–973.

na

12.

13. Wu C, Chang J. A review of bioactive silicate ceramics. Biomed mater. 2013; 8: 032001. 14. Zhai W, Lu H, Wu C, Chen Li, Lin X, Naoki, Kawazoe, Chen G, Chang J. Stimulatory effects of the ionic

ur

products from Ca-Mg-Si bioceramics on both osteogenesis and angiogenesis in vitro. Acta Biomater. 2013; 9: 8004-14.5.

Jo

15. Diba M, Goudouri O-M, Tapia F, Boccaccini AR. Magnesium-containing bioactive polycrystalline silicatebased ceramics and glass-ceramics for biomedical applications. Curr Opin Solid State Mater Sci. 2014; 18: 147167.

16. Zhai W, Lu H, Lei Chen L, Lin X, Huang Y, Dai K, Naoki K, Chen G, Chang J. Silicate bioceramics induce angiogenesis during bone regeneration. Acta Biomater. 2012; 8: 341-349. 17. Wu C, Chang J. Degradation, bioactivity, and cytocompatibility of diopside, akermanite, and bredigite ceramics. J Biomed Mater Res Part B: Appl Biomater. 2007; 83B: 153-60. 18. Xu S, Lin K, Wang Z, Chang J, Wang L, Lu J, Ning C. Reconstruction of calvarial defect of rabbits using porous calcium silicate bioactive ceramics. Biomaterials. 2008; 29: 2588-96. 15

19. Wu C, Ramaswamy Y, Zreiqat H. Porous diopside (CaMgSi2O6) scaffold: A promising bioactive material for bone tissue engineering. Acta Biomater. 2010; 6: 2237-45. 20. Sainz MA, Pena P, Serena S, Caballero A. Influence of design on bioactivity of novel CaSiO 3-CaMg(SiO3)2 bioceramics: In vitro simulated body fluid test and thermodynamic simulation. Acta Biomater. 2010; 6: 2797807. 21. Xie J, Yang X, Shao H, Ye J, He Y, Fu J, Gao C, Gou Z. Simultaneous mechanical property and biodegradation improvement of wollastonite bioceramic through magnesium dilute doping. J Mech Behav Biomed. 2016; 54: 60-71. 22. Xie J, Shao H, He D, Yang X, Yao C, Ye Y, Fu J, Gou Z. Ultrahigh strength of three-dimensional printed diluted magnesium doping wollastonite porous scaffolds. MRS Commun. 2015; 5(04): 631-639. 23. He D, Zhuang C, Xu S, Ke X, Yang X, Zhang L, Yang G, Chen X, Mou X, Liu A, Gou Z. 3D printing of Mg-

ro of

substituted wollastonite reinforcing diopside porous bioceramics with enhanced mechanical and biological performances. Bioactive Mater. 2016; 1: 85-92.

24. Chen Z, Li Z, Li J, Li J, Liu C, Lao C, Fu Y, Liu C, Li Y, Wang P, He Y. 3D printing of ceramics: A review. J Eur Ceram Soc. 2019; 39: 661-87.

25. Roleček J, Pejchalová L, Martínez-Vázquez FJ, Miranda González P, Salamon D. Bioceramic scaffolds

1595-602.

-p

fabrication: Indirect 3D printing combined with ice-templating vs. robocasting. J Eur Ceram Soc. 2019; 39: 26. Schmidt J, Elsayed H, Bernardo E, Colombo P. Digital light processing of wollastonite-diopside glass-ceramic

re

complex structures. J Eur Ceram Soc. 2018; 38: 4580-4.

27. Bose S, Ke D, Sahasrabudhe H, Bandyopadhyay A. Additive manufacturing of biomaterials. Prog Mater Sci.

lP

2018; 93: 45-111.

28. Deville S, Saiz E, Tomsia AP. Freeze casting of hydroxyapatite scaffolds for bone tissue engineering. Biomaterials. 2006; 27: 5480-5489.

29. Bai F, Wang Z, Lu J, Liu J, Chen G, Lv R, Wang J, Lin K, Zhang J, Huang X. The correlation between the

na

internal structure and vascularization of controllable porous bioceramic materials in vivo: a quantitative study. Tissue Eng Part A. 2010; 16: 3791-803.

30. Ma H, Feng C, Chang J, Wu C. 3D-printed bioceramic scaffolds: From bone tissue engineering to tumor

ur

therapy. Acta Biomater. 2018; 79: 37-59.

31. Shao H, Yang X, He Y, Fu J, Liu L, Ma L, Zhang L, Yang G, Gao C, and Gou Z. Bioactive glass-reinforced

Jo

bioceramic ink writing scaffolds: Sintering, microstructure and mechanical behavior. Biofabrication. 2015; 7: 035010.

32.

Liu L, He D, Yang X, Zhang L, Yang G, Dai X, Gou Z. Design strategy for the multi-layer core-shell

bioceramics with controlled chemistry. J Am Ceram Soc. 2016; 99: 2243-2252.

33. Zhuang C, Ke X, Jin Z, Zhang L, Yang X, Xu S, Xie L, Edem Prince G-A, Pan Z, Gou Z. Core-shell-structured nonstoichiometric bioceramic microspheres for improving osteogenic capability. J Mater Chem B. 2017; 5: 8944-8956.

16

34. Fu J, Zhuang C, Qiu J, Ke X, Yang X, Jin Z, Zhang L, Yang G, Xie L, Xu S, Gao C, Gou Z. Core-shell biphasic microspheres with tunable density of shell micropores providing tailorable bone regeneration. Tissue Eng. Part A. 2019; 25: 588-602. 35. Hayashi S, Okada K, N. ōtsuka. Sintering behaviour of diopside, CaMgSi2O6, from various powder preparation methods. J Mater Sci Lett. 1990; 9: 382-385. 36. Kokubo T, Takadama H. How useful is SBF in predicting in vivo bone bioactivity? Biomaterials. 2006; 27: 2907-15. 37. Hench LL. Bioceramics, J Am Ceram Soc. 1998; 81: 1705-1728. 38. Dorozhkin SV. Calcium orthophosphate bioceramics. Ceram Int. 2015; 41: 13913-13966. 39. Baino F, Novajra G, Vitale-Brovarone C. Bioceramics and Scaffolds: A Winning Combination for Tissue Engineering. Front Bioeng Biotechnol. 2015; 3: 202.

ro of

40. Wu C, Luo Y, Cuniberti G, Xiao Y, Gelinsky M. Three-dimensional printing of hierarchical and tough mesoporous bioactive glass scaffolds with a controllable pore architecture, excellent mechanical strength and mineralization ability. Acta Biomater. 2011; 7: 2644-50.

41. Shao H, Ke X, Liu A, Sun M, He Y, Yang X, Fu J, Liu Y, Zhang L, Yang G, Xu S, Gou Z. Bone regeneration in 3D printing bioactive ceramic scaffolds with improved tissue/material interface pore architecture in thin-wall

-p

bone defect. Biofabrication. 2017; 9: 25003.

42. Jin Z, Wu R, Shen J, Yang X, Shen M, Zhang L, Yang G, Gao C, Gou Z, Xu S. Nonstoichiometric wollastonite bioceramic scaffolds with core-shell pore struts and adjustable mechanical and biodegradable properties. J

re

Mechan Behav Biomed Mater. 2018; 88: 140-149.

43. Paredes C, Martínez-Vázquez FJ, Pajares A, Miranda P. Development by robocasting and mechanical

4375-4383.

lP

characterization of hybrid HA/PCL coaxial scaffolds for biomedical applications. J Euro Ceram Soc. 2019; 39: 44. Zhang W, Feng C, Yang G, Li G, Ding X, Wang S, Dou Y, Zhang Z, Chang J, Wu C, Jiang X. 3D-printed scaffolds with synergistic effect of hollow-pipe structure and bioactive ions for vascularized bone regeneration.

na

Biomaterials 2017; 135: 85-95.

45. Feng C, Zhang W, Deng C, Li G, Chang J, Zhang Z, Jiang X, Wu C. 3D Printing of lotus root-like biomimetic materials for cell delivery and tissue regeneration. Adv Sci. 2017; 4: 1700401.

ur

46. Li T, Zhai D, Ma B, Xue J, Zhao P, Chang J, Gelinsky M, Wu C. 3D Printing of hot dog-like biomaterials with hierarchical architecture and distinct bioactivity. Adv Sci. 2019, 6: 1901146.

Jo

47. Nonami T, Tsutsumi S. Study of diopside ceramics for biomaterials. J Mater Sci Mater Med .1999; 10: 475-479. 48. Kobayashi T, Okada K, Kuroda T, Sato K. Osteogenic cell cytotoxicity and biomechanical strength of the new ceramic Diopside. J Biomed Mater Res. 1997; 37:100-7.

49. Champion E. Sintering of calcium phosphate bioceramics. Acta Biomater. 2013; 9: 5855-75. 50. Descamps M, Boilet L, Moreau G, Tricoteaux A, Lu J, Leriche A, Lardot V, Cambier F. Processing and properties of biphasic calcium phosphates bioceramics obtained by pressureless sintering and hot isostatic pressing. J Eur Ceram Soc. 2013; 33: 1263-70. 51. Wu C, Fan W, Zhou Y, Luo Y, Gelinsky M, Chang J, Xiao Y. 3D-printing of highly uniform CaSiO3 ceramic scaffolds: preparation, characterization and in vivo osteogenesis. J Mater Chem. 2012; 22: 12288. 17

52. Wopenka B, Pasteris JD. A mineralogical perspective on the apatite in bone. Mater Sci Eng C. 2005; 25: 13143. 53. Hugo R. Fernandes, Anuraag Gaddam, Avito Rebelo, Daniela Brazete, George E. Stan, José M. F. Ferreira. Bioactive glasses and glass-ceramics for healthcare applications in bone regeneration and tissue engineering.

Jo

ur

na

lP

re

-p

ro of

Materials (Basel). 2018; 11: 2530.

18

Figure Captions Figure 1. XRD patterns of CSi and Dio powders (A) and the magnified peak at 29.7o/2θ representing the (-2, 2, 1) crystal plane when substituted by Zn and Sr (B).

re

-p

ro of

Figure 2. SEM images of the fracture microstructures of as-sintered bioceramic scaffolds.

lP

Figure 3. SEM images of core-shell pore strut of bioceramic scaffolds. The images in (A1) and (B1) representing the high-magnification of (A) and (B), and the EDX spectra showing the element composition of the core and

Jo

ur

na

shell layer of CSi@Dio and CSi@Dio-p scaffolds, respectively. Bar: 100 μm.

19

ro of

Figure 4. Compressive strength (A) and Young’s Modulus (B) and representative stress-strain curves (C-F) of the

na

lP

re

-p

bioceramic scaffolds after sintering at 1250℃.

Figure 5. SEM images of the evolution of the core-shell structure (A) with a 1000 magnification of representative CSi@Dio and CSi@Dio-p scaffolds (B); weight loss and mechanical decay (C,D) of the scaffolds immersing

Jo

ur

in Tris-HCl buffers for 8 weeks.

20

Figure 6. Bioactive ion release of the scaffolds of Ca (A), Mg (B), Si (C), Zn (D), and Sr (E) soaking in Tris buffer for 4 h‒14 d.

Figure 7. SEM images of the microstructures of surface (A-F) and factures (G-L) in the pore struts for the CSi@Dio

ro of

and CSi@Dio-p scaffolds after immersion in Tris buffer for different time stages. Dotted circle and arrow showing the surface microspores in the Dio shell layer, and dotted line displaying the core-shell interface

ur

na

lP

re

-p

of pore struts.

Jo

Figure 8. SEM images of the surface morphology of the bioceramic scaffolds after immersion in SBF for 7 days. Inset showing the macroporous structures of the bioceramic scaffolds by SEM observation and the EDX spectra in the surface layer after soaking in SBF.

21

ro of

Scheme 1. Schematic illustration for 3D printing calcium silicate@diopisde (CSi@Dio) bioceramic scaffolds with

Jo

ur

na

lP

re

-p

core-shell-structured nozzle.

22

Tables

Table 1. ICP of 4% Sr or Zn doping Dio Ca Mg （ppm） （ppm）

Si

Sr

（ppm） （ppm）

Zn（ppm） Sr/(Sr+Mg) Sr/(Sr+Ca) Zn/(Zn+Ca) Zn/(Zn+Mg)

12.63

14.24

29.4

0.95

0.06

1.79%

3.31%

╱

╱

4%Zn-Dio

9.21

7.6

18.01

0.05

0.89

╱

╱

0.25%

4.78%

Jo

ur

na

lP

re

-p

ro of

4%Sr-Dio

23

Table 2. Element composition (EDX analysis, atom%) in the pore-strut core or shell layer of CSi@Dio and CSi@Dio-p scaffolds. CSi@Dio

CSi@Dio

CSi@Dio-p

CSi@Dio-p

O

62.33

63.75

63.90

60.73

Mg

0.75

8.15

1.15

7.83

Si

17.93

20.68

17.56

19.13

Ca

18.98

7.41

17.39

12.31

Ca/Mg

25.31

0.91

15.12

1.57

Jo

ur

na

lP

re

-p

ro of

Elements

24