Destructive and non-destructive behavior of nickel

Destructive and non-destructive behavior of nickel oxide doped bioactive glass and glass-ceramic Vikash Kumar Vyas, Arepalli Sampath Kumar, S. P. Singh & Ram Pyare

Journal of the Australian Ceramic Society ISSN 2510-1560 Volume 53 Number 2 J Aust Ceram Soc (2017) 53:939-951 DOI 10.1007/s41779-017-0110-2

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Author's personal copy J Aust Ceram Soc (2017) 53:939–951 DOI 10.1007/s41779-017-0110-2

RESEARCH

Destructive and non-destructive behavior of nickel oxide doped bioactive glass and glass-ceramic Vikash Kumar Vyas 1 & Arepalli Sampath Kumar 1 & S. P. Singh 1 & Ram Pyare 1

Received: 24 March 2017 / Revised: 27 June 2017 / Accepted: 29 June 2017 / Published online: 18 July 2017 # Australian Ceramic Society 2017

Abstract Nickel oxide substituted bioactive glasses (45S5) have been prepared by melting and annealing techniques. The doping of Ni2+ ion from 0 to 1.65 mol% of NiO was done to replace Si4+ ion and yield a charge balanced (CB) bioactive glass. The Ni2+ ion would enter into [SiO4]4− network as [NiO4]2− tetrahedra due similar charge/size ratio, but depending upon oxygen environment, it may act as modifier also in octahedral coordination in the glass. Polycrystalline bioactive glass-ceramics were prepared through controlled heat treatment. The glass and glass-ceramic structure was evaluated using FTIR and XRD techniques. The crystalline phases in bioactive glass-ceramics were identified using X-ray difractometry. The SEM micrographs of the samples after chemical treatment with simulated body fluid (SBF) for definite time of 15 days had shown the formation of hydroxyl carbonate apatite (HCP) layer on their surface which indicated that NiO had no opposite effect on the overall bioactivity. The destructive tests like microhardness, compressive, flexural strengths, and the non-destructive tests of elastic moduli were carried out. Both the results indicated that substitution of nickel oxide by silica in 45S5 bioactive glass and glass-ceramic influenced the structure and enhanced its density, compressive, flexural strength, micro hardness, and elastic properties.

Keywords Bioactive glasses . Bioactive glass-ceramics . Mechanical properties . Elastic modulus

* Vikash Kumar Vyas [email protected]; [email protected]

1

Department of Ceramic Engineering, Indian Institute of Technology (Banaras Hindu University), Varanasi 221005, India

Introduction Bioactive glasses, first developed by L.L Hench in 1969, are a class of biomaterials that have ability to bond to the soft and hard tissues of bone. The wt% composition of Hench bioglass® (45S5) is 45SiO2-24.5CaO-24.5Na2O6P2O5 [1]. The bioactive glasses and its derivatives are commonly used material in clinical application, such as ossicular implantation for alleviating conductive hearing loss and dental applications [2–8]. One of the major applications of bioactive glass and glass-ceramic is as an artificial bone graft. Therefore, it is a promising material in the field of biomedical application. It has inferior mechanical properties in comparison to cortical bone. However, it suffers from a mechanical weakness and low fracture toughness due to an amorphous nature of glass and it may not be suitable for load-bearing applications [9]. The mechanical strength of a glass depends on its elastic properties, shape, surface/internal defects and the character of the force to which it is subjected. These factors inevitably degrade the mechanical strength of a material [10]. Thus, one cannot obtain the intrinsic strength by using destructive mechanical test on a larger specimen. The ultrasonic non-destructive testing has been found to be one of the best techniques to study the elastic properties, microstructure, and mechanical behavior as well as to evaluate elastic constants of glasses [11–15]. The investigation of elastic properties of bioactive glasses has stimulated many researchers [11, 16, 17], and significant information about the same has been obtained. The elastic moduli which may in turn be measured by ultrasonic velocities have the influence on physical parameters. The structural configuration between glass former and modifier of compositions of glasses has got its significant effects on ultrasonic velocities [18, 19].

Author's personal copy 940

Shrivastava et al. [20] also investigated on the elastic properties of 45S5 bioglass® substituted with Cu2+ ion for Si4+ in bioactive glass and glass-ceramics and reported the enhancement in physico-mechanical properties. Smith et al. [21] studied the structural characterization of hypoxia-mimicking 45S5 glass and found that nickel adopted a mixed structural role in bioactive glass occupying both network forming tetrahedral and network modifying fivefold symmetries. Structural correlation between 45S5 bioactive glass and nickel oxide doped bioactive glass had revealed that isomorphic substitution of nickel oxide in the silicate network in tetrahedral coordination or modifier in higher coordination did not adversely affect the existing glass structure. Sampth et al. [22] replaced silica by barium oxide and reported an increase in density and flexural strength of the bioactive glass due to replacement of Si4+ ion with a heavier Ba2+ ion on increasing BaO content in the glass. Tripathi et al. [23] found an increase in elastic moduli in Li2O–CaO–Al2O3–P2O5–SiO2 bioactive glasses with increasing Al2O3/Li2O ratio up to 2.0 mol% of Al2O3 and then after that became stable with increasing the ratio in the glass. Gaafar et al. [24] have also earlier mentioned that the change in behavior of Poisson’s ratio in cobalt oxide doped borate glasses was due to change in the type of bonding in the glass structure. The role of nickel oxide as a colorant and decolorizer in smaller quantities present in traditional glasses is well known. Various systematic investigations have been carried out by earlier workers [25–39] on optical, redox (Ni/Ni2+), and magnetic properties of nickel-contained glasses with small concentrations of NiO. It is agreed that nickel oxide in minor quantities present in glasses has caused little influence on the glass system, but whenever it is present as major ingredient in the glass composition, it has got its considerable influence on the structure of the glasses [28, 34–36]. Singh and Singh [34] studied the thermodynamic activity of nickel oxide in alkali nickel silicate glasses (R2O·NiO·SiO2 where R+═Li+, Na+, and K+ ions), with varying concentrations of NiO in pure nickel crucible at high temperature in air as furnace atmosphere. After equilibrating the Ni 0 /Ni 2+ redox in the melts, they established the relation between logarithm of activity coefficient (log γNiO) and concentration of NiO in the glass. They found that Henry’s law was obeyed in the lower concentration range, whereas the curve deviated from this in the higher concentration range of NiO in a particular silicate glass at a fixed temperature. This indicated that the addition of NiO in the glass in major amount had caused structural changes in the glassy

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matrix. Paul [31] had also investigated the activity of NiO at a fixed temperature at 1400° in air atmosphere in alkali borate glasses as well as with varying temperatures (850–1100 °C) and partial pressure of oxygen (pO 2 = 10 − 9 –10 − 1 4 atm) in 20Na 2 O·80B 2 O 3 and 20K2O·80B2O3 glass. The authors inferred that the solubility of NiO increased with an increase in alkali oxides content and attained a maxima up to certain limits of alkali concentration and then further it decreased with increasing alkali contents in the glass. He mentioned that the low solubility of NiO in high alkali borate glasses was due to weak donor capacity of [BO3/2] group and high O2− ion activity in the molten glass. Pretorius and Muan [32] also studied the activity of NiO in molten silicate glasses at high temperature on varying pO2 and observed a decrease in the activity coefficient of NiO with increasing basicity of the glass for non-bridging oxygen/silicon ratio [O−/Si] > 1.5 as well as an increase with decreasing basicity with [O−/ Si] ratio < 1.0. During their studies, Henry’s law was also found to obey for a relation between log γNiO versus concentration of NiO for the glasses of similar basicity. Similar results were also observed regarding effect of additions of oxides of copper, iron, and vanadium in major quantities on the structure of alkali–copper–silicate, iron-doped soda–lime–silica, and sodium borovanadate glasses by earlier workers [36–38]. Suresh et al. [39] have also investigated the spectroscopic features of lead–bismuth–silicate glasses containing 0–1.0 mol% NiO and observed that Ni2+ ions were present as both in octahedral and tetrahedral symmetries in the glass. The authors mentioned that in increasing the concentration of Ni2+ ion resulted an increase in its octahedral coordination which also dominated the tetrahedral coordination of Ni2+ ion in the glass. Their IR, luminescence, and Raman spectroscopic studies suggested that increasing randomness in the glassy matrix was due to an increase in NiO concentration which acted as modifier in the glass and depolymerized the host glass network. Bioactive glasses containing NiO have offered an existing route for potential delivery system of Ni2+ ions with tissue regeneration scaffolds due to its ability to incorporate a large variety of elements such as Ca and P and their appreciably controlled dissolution properties within physiological fluids. Ni2+ ions are being considered as a possible alternative to growth factors and genetic approaches in tissue engineering because of their biocompatibility, easy processing, and tunable release kinetics. In view of the important role played by Ni2+ ion in tissue engineering, the present work was undertaken.

Author's personal copy J Aust Ceram Soc (2017) 53:939–951 Table 1

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Composition of bioactive glasses (mol %)

Sample

SiO2

Na2O

CaO

P2O5

NiO

45S5

46.14

24.40

26.91

2.55

0.00

NiO-1

45.64

24.41

26.94

2.60

0.41

NiO-2 NiO-3

45.20 44.71

24.44 24.46

26.96 26.98

2.61 2.61

0.82 1.23

NiO-4

44.25

24.49

27.01

2.61

1.65

Experimental procedures Selecting composition and glass preparation The bioactive glass composition was formulated from Na2O–CaO–SiO2–P2O5 glass system. First, the bioactive 45S5 glass, having mol % composition [46.1SiO 2 24.4Na2O-26.9CaO-2.5P2O5], was prepared. Then the proposed bioglass® containing chemical composition (46.1 − X) SiO 2 -24.4Na 2 O-26.9CaO-2.5P 2 O 5 (where X = 0–1.65% of NiO) was further prepared. In this study, the mol percent of CaO, Na2O, and P2O5 was kept constant and SiO 2 was partially replaced with NiO. The compositions of bioactive glass® are given in Table 1. In the bioactive glass® of the system, 45S5 we have added mol% nickel oxide (0– 1.65 mol%) for SiO2 and prepared using normal melting and annealing techniques. Nickel oxide was substituted in the glass composition for silica in different concentrations to yield a charge balanced (CB) series of bioglass® as less SiO2 is associated with more NBO [−O−] based on its dual role in the glass. The other components of the bioactive glass® are kept constant. The materials used include fine-grained quartz for silica. Sodium and calcium carbonate were introduced in the form of their respective anhydrous carbonates. P2O5 was introduced in the form of ammonium dihydrogen orthophosphate [NH 4 H 2 PO 4 ]. The weighed batches were mixed thoroughly for 45 min and melted in pure (99.9%) alumina crucibles of 100 ml capacity. The glass samples were from different batches. Each of the 100 g glass batches for five bioglass samples containing 0– 1.65% NiO was prepared and melted separately in 100 ml alumina crucible in identical conditions. It is a fact when silicate glasses were melted previously in alumina crucibles for a prolonged period of 70 h at 1400 °C, an appreciable amount of alumina was dissolved in the glass melts and influenced its composition. However, it has been earlier observed that alumina does not dissolve in silicate and phosphate glasses [36]

considerably into glass melts during small duration of 6 h of melting at high temperature [40, 41]. On plotting the portion of iron present in ferrous state (Fe2+/ Fe) against the duration of heat treatment, Baak and Hornyak [40] found that there was a smooth curve for the experiments conducted in an alumina crucible where as a zig-zag relationship was obtained for heat treatments carried out at high temperature in a platinum crucible. They inferred that platinum interfered with the Fe2+/Fe3+ redox equilibrium, but alumina did not interfere with the same in the sodium disilicate glass. Majhi et al. [41] also reported in calcium–iron–phosphate glasses by gravimetric analysis that the glass melts did not contaminate with Al2O3 during melting in alumina crucibles at 1200 °C for a period of 6 h in air as furnace atmosphere. The XRD patterns (Fig. 3) of the bioactive glassceramic derivatives in the present study have also not shown any separate crystalline phase containing alumina in the samples. The melting was carried out in an electric furnace at 1400 ± 10 °C for 3 h in air as furnace atmosphere, and homogenized melts were poured on preheated aluminum sheet. The prepared bioactive glass samples were directly transferred to a regulated muffle furnace at 490 °C for annealing. After 2 h of annealing, the muffle furnace was cooled to room temperature with controlled rate of cooling at 10 °C per min.

Fig. 1 DTA curve of bioactive glasses


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Table 2 Heat treatment temperatures used for nucleation and crystal growth of bio-glasses Sample

Nucleation (°C)

Time (h)

Crystallization (°C)

Time (h)

45S5

506

6

680

3

NiO-1

463

6

630

3

NiO-2 NiO-3

452 450

6 6

614 606

3 3

NiO-4

447

6

609

3

Differential thermal analysis measurements Differential thermal analysis (DTA) is used to find out the nucleation and crystallization temperatures of the nickel oxide substituted bioactive glass and base glass. Fine powder of bioactive glasses was made using an agate mortar and pestle and analyzed using a differential thermal analyzer (SETARAM, France) at a heating rate of 10 °C per minute under a stream of argon atmosphere using alumina as a reference material. This experiment was carried out in the temperature range of 50 to 800 °C. Heat treatment regime (conversion of glass to glass-ceramic) Based on DTA results (Fig. 1), each bioactive glass sample was firstly heated slowly to the nucleation temperature (Tg) from room temperature to 600 °C for the formation of sufficient nuclei sites, and then after holding it for a definite time of 6 h, it was further heated to reach the second chosen crystal growth temperature (Tc = 600–800 °C) for performing the perfect crystal growth process. After a second hold for the specific time period of 3 h, the sample was left inside the muffle furnace for cooling to room temperature at a rate of 10 °C/h. The temperatures are given in Table 2 for each bioactive glass sample. Bioactive glass and glass-ceramic characterization In order to identity the crystalline phase present in the heat-treated bioactive glass samples, the glass-ceramic samples were ground to 75 μm and the fine powders were subjected to X-ray diffraction analysis (XRD). A RIGAKU-Miniflex II diffractometer adopted Cu-Kα radiation (λ = 1.5405A°) with a tube voltage of 40 KV and a current of 35 mA in a 2θ range between 10° and 80°. The step, size, and measuring speed was set to 0.02° and 1° per min respectively in the present investigation. The JCPDS-International Centre for diffraction

Data Cards were used as a reference to find out the different phases formed. The structure of bioactive glass samples was measured at room temperatures in the frequency range of 4000–400 cm−1 using a Fourier transform infrared spectrometer (Bruker Tensor 27 FTIR, USA). The fine bioactive glass and glass-ceramic powder samples were mixed with KBr in the ratio of 1:100, and the mixtures were subjected to an evocable die at load of 12 bar pressure to produce clear standardized discs. The prepared discs were immediately subjected to IR spectrometer to measure the reflectance spectra in order to avoid moisture attack. Mechanical properties The densities of the ground and polished bioactive glass and glass-ceramic samples were measured by Archimedes’s principle at room temperature. Three pieces of each composition were taken for measuring the mechanical properties of the samples. Compressive strength of the base glass, nickel oxide substituted bioactive glasses, and their respective glassceramics (2 × 2 × 1 cm3 size) according to ASTM D3171 were subjected to compression test. The test was performed using Instron Universal Testing Machine at room temperature (cross speed of 0.05 cm/min and full scale of 5000 kgf). Flexural strength was determined according to ASTM Standard: C158–02(2012). Polished bioactive glass samples are used for hardness testing, and the size of sample was 1.2 cm × 1.2 cm × 1.0 cm according to ASTM Standard: C730-98. The indentations have been made for loads ranging between 30 and 2000 N, applied at a velocity of 1 mm/s, and allowed to equilibrate for 16 s before measurement. Elastic properties of glass and glass-ceramics The ultrasonic wave velocities (longitudinal and shear) for nickel oxide substituted bioglass® and base glass and its ceramic derivates were measured using the Olympus instrument (M-45, USA) made in the USA. Bioactive glass and glass-ceramic samples were cut and polished in cubic pieces, and the couplant glycerin was used for finding longitudinal velocities and sonfech shear gel for the shear velocities of bioactive glass and its ceramic derivative. Using the appropriate formula, the Young’s, shear, bulk modulus of elasticity, and Poisson’s Ratio were found. Assessment of bioactivity by SEM In order to ensure that the addition of Ni2+ ion would not have any adverse effect on the overall bioactivity, the scanning electron microscopic (SEM) study of the bioactive glass and

Author's personal copy J Aust Ceram Soc (2017) 53:939–951

Fig. 2 XRD pattern of bioactive glasses

glass-ceramic samples was done before and after SBF treatment for definite period of 15 days. The surface morphology of samples was analyzed before and after SBF treatment using a scanning electron microscope (SEM—Inspect S50, FEI). The SBF was prepared with analytical reagent grade chemicals by Kokubo’s method [42]. The samples were coated with gold (Au) by sputter coating instrument before analyzing by SEM.

Results and discussion Differential thermal analysis of bioactive glasses The DTA plots of base and nickel-substituted bioactive glasses are presented in Fig. 1. The DTA patterns of all the nickel substituted bioactive glass® samples (NiO-1, NiO-2 NiO-3, and NiO-4) are similar to each other with respect to 45S5 bioactive glass® overall range of temperature. The Tg value for the bioactive glass® sample no NiO-1 has been clearly determined from its DTA curve, as tabulated in Table 2, along with Tg of other samples. The DTA curves have shown only one exothermic peak corresponding to crystallization of a single phase in the glass-ceramic samples, and the DTA curve shows the decrease in exothermic as well as endothermic peaks. This decrease in the peaks is due to incorporation of nickel oxide which entered the glass network and formed Si–O–Ni bonds which replaced the stronger Si–O–Si bonds. Thus, the glass network becomes weaker, and Tg and Tc decreased [43]. Further, a decrease in the glass transition

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Fig. 3 XRD pattern of glass-ceramics

temperature (Tg) in DTA curves (Fig. 1), with increasing addition of NiO up to 2.0 mol% in a concentration-dependent manner, had shown that Ni2+ ion acted as an intermediate in smaller quantities. Because of the similar ionic radii and therefore similar charge to size ratio, it might be expected that Ni2+ (0.58 Å) ion would adopt similar structural role as Co2+ (0.55 Å) ion which had been shown earlier to act as a network intermediate [21]. It is mentioned herewith that due to charge/size ratio of Ni2+ ion, it would enter into the [SiO4]4− network as [NiO4]2− species in tetrahedral coordination but it could act also as network modifier in octahedral coordination depending upon the availability of oxygen environment. The concentration of NiO in the present bioactive glass system is small and only up to 0–1.65 mol%. During the small addition of NiO for SiO2 in the initial stages, the heavier Ni2+ ion might have entered as [NiO4]2− species and replaced lighter Si4+ ion in [SiO4]4− network until it remained as is, which has resulted the formation of Ni–O– Si bonds in silicate structure. It is supported in the preceding discussion that Smith et al. [21] investigated the structural characterization of bioactive glasses doped with 4.0 mol% NiO using neutron diffraction technique and reported the Ni–O–Si bonding in silicate network with one third of Ni2+ ion occupying a network forming tetrahedral geometry and its two-third occupying a fivefold coordination in the bioglass. Smith et al. [21] had also confirmed, by isomorphic substitution, that Ni2+ and Co2+ ions are both isostructural in the bioactive glasses.


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Fig. 6 Variation in density of bioactive glasses and glass-ceramics with NiO concentration

Fig. 4 FTIR reflectance spectra of bioactive glasses

The present trend of a decrease in Tg with increasing concentration of NiO up to 1.65 mol% in our bioactive glass samples is well supported with earlier observations made by Azevedo et al. [43] for their bioactive glasses. They have mentioned that a decrease in glass transition temperature determined by differential scanning calorimetry (DSC) had

shown that the addition of CoO in smaller amounts up to 2.0 mol% acted as an intermediate and further of addition up to 4.0 mol% CoO in larger quantities had acted as a modifier in their glasses. According to Zachariasen’s rules [44, 45] and Dietzel criteria [46] for glass formation, Co2+ ion may also act as an intermediate oxide due to its relatively high charge-tosize ratio like Mg2+ and Ca2+ ions. It would be worth to mention herewith that it is the coordination number (CN) and charge/size ratio of Co2+ ion which had determined its dual role as an intermediate and modifier in the glass. Therefore, accordingly, Co2+ ion in higher coordination can act as a modifier which replaces bridging oxygen (−O−) with non-bridging oxygen (−O−) in the bioglass.

X-ray diffraction patterns of bioactive glass and glass-ceramic

Fig. 5 FTIR reflectance spectra of bioactive glass-ceramics

The XRD patterns of glass and glass-ceramic for base glass and nickel oxide substituted glasses are presented in Figs. 2 and 3, respectively. The observed results indicate that the glasses have amorphous structure, and there is no indentation for the presence of crystalline phases. Further, Fig. 2 clearly shows that the nickel oxide is completely dissolved in the glass matrix. The XRD patterns for base glass-ceramic and nickel oxide substituted glass-ceramic are shown in Fig. 3. The XRD patterns of the glass-ceramic substituted and unsubstituted show the presence of crystalline phase of sodium calcium silicate [Na2Ca2Si3O9] (card no.: PDF#01-1078& PDF#02-0961,Na2CaSi3O8 (card number: PDF#12–0671) [20]. It can be clearly seen from the XRD results that the presence of nickel oxide in the glass-ceramic did not show


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Fig. 7 Variation in compressive strength of glasses and glass-ceramics with NiO concentration

Fig. 9 Variation in microhardness of glasses and glass-ceramics with NiO concentration

nickel oxide as separate crystalline phase. This may be related to their relatively low content in bioactive glass compositions.

ceramic (NiO-1C) shows the peaks at 458, 691, 1045, 1485, and 3779 cm−1 .The resultant IR spectra at 495 and 458 cm−1 are associated with a Si–O–Si symmetric bending mode, and the band at 716 and 691 cm−1 corresponds to the Si–O–Si symmetric stretch of non-bridging oxygen atoms between tetrahedral. It was observed that the intensity of the band was recreated as the nickel oxide substitution in the base glass; therefore, the nickel oxide increases the non-bridging oxygen in the network. The major band at about 1007 and 1045 cm−1 is attributed to Si–O–Si stretching. The small band at 1485 and 1485 cm−1 is attributed to C–C vibration mode. It was observed that the intensity of the IR peak increased as the concentration of nickel oxide increase which is due to the breaking of Si–O–Si network. The small broad band centered at about 3768 and 3779 cm−1 can be assigned to hydroxyl group (−OH) which may be due to the presence of adsorbed water molecules. Figures 4 and 5 depict the infrared frequencies and related functional structural groups in the bioactive

FTIR analysis of bioactive glass and glass-ceramic Figures 4 and 5 shows the FTIR reflection spectra of the base and nickel oxide substituted bioactive glass and glass-ceramic. All the bioactive glass and glass-ceramic samples are showing similar trend behavior. FTIR reflection spectra bands of all the glasses confirm the main characteristic of silicate network, and this may be due to the presence of SiO2 as a major constituent. Therefore, the bioactive glass (NiO-1) shows the peaks at 495, 716, 1007, 1485, and 3768 cm−1 and glass-

Table 3 Density, longitudinal, and shear velocities of bioactive glass and glass-ceramic Glasses

Glass-ceramic

NiO-doped bioactive glass

NiO-doped glass ceramic

Sample ρ(g/cm3) VL(m/s) VT (m/s) ρ(g/cm3) VL(m/s) VT (m/s)

Fig. 8 Variation in flexural strength of glasses and glass-ceramics with NiO concentration

45S5 NiO-1 NiO-2 NiO-3 NiO-4

2.77 2.82 2.83 2.86 2.89

5908 5926 5936 5949 5967

3356 3378 3389 3406 3419

2.913 2.918 2.924 2.949 2.951

6514 6524 6538 6551 6573

3732 3749 3764 3775 3793

Author's personal copy 946 Table 4 Sample

J Aust Ceram Soc (2017) 53:939–951 Young’s modulus, shear modulus, bulk modulus, and Poisson’s ratio of glasses and glass-ceramics Glasses

Glass-ceramic

E(Gpa) G(Gpa) K(Gpa) NiO doped 45S5 bioactive glasses

σ

E(Gpa) G(Gpa) NiO doped glass-ceramic

K (Gpa)

σ

45S5

78.86

31.19

54.99

0.261

102.02

40.57

69.40

0.255

NiO-1

81.10

32.17

56.29

0.259

102.91

41.01

69.16

0.253

NiO-2 NiO-3

81.53 83.54

32.51 33.17

56.15 56.83

0.258 0.255

103.76 105.46

41.42 42.02

69.73 70.31

0.252 0.250

NiO-4

85.10

33.78

57.65

0.254

106.44

42.45

70.68

0.249

glass and glass-ceramic [47]. The bioactive glasses substituted with nickel oxide do not show noticeable changes in the IR spectra bands. Density, compressive, flexural strength, and microhardness of glass and glass-ceramic Figure 6 shows the density of nickel oxide substituted 45S5 bioactive glass and glass-ceramic which was measured by Archimedes’ principle. It was observed that the densities of the samples increased with increasing nickel oxide content from 2.77 to 2.89 g/cm3 and 2.91 to 2.95 g/cm3 for glasses and glass-ceramics, respectively. It may be attributed to partial replacement of lighter SiO2 (2.64 g/cm3) with heavier NiO (6.67 g/cm3) in the composition. The compressive strengths of the glass and glass-ceramic were determined and reported as 68.92, 76.92, 78.99, 79.98, and 82.13 MPa for sample no. 45S5, NiO-1, NiO-2, NiO-3, and NiO-4 and in glass-ceramic 113, 118, 121, 126, and 121 MPa, respectively, as given in Fig. 7. The flexural strength of these samples were also measured as 42.43, 55.15, 56.42, 60.52, and 66.58 MPa for

Fig. 10 Variation in Young’s Modulus of glasses and glass-ceramics with NiO concentration

samples 45S5, NiO-1, NiO-2, NiO-3, and NiO-4 and for glass-ceramics 104, 108, 110, 113, and 121 MPa as given in Fig. 8 and micro hardness for glass 5.35, 5.48, 5.51, 5.77, and 5.99 as well as glass-ceramic 7.68, 7.81, 7.86, 8.06, 8, and 8.14 as presented in Fig. 9. It can be easily understood that bioactive glass and their respective glass-ceramic substituted with nickel oxide have demonstrated comparatively higher compressive and flexural strengths and microhardness, respectively. This is also in good agreement with density values as presented in Fig. 6. This increase may be due to the fact that Ni2+ ion might have acted as network intermediate, thus leading to more compactness of the glass structure [48]. By isotopic and isomorphic substitution of nickel, Smith et al. [21] had shown that there were no appreciable differences in structural correlation between 45S5 bioactive glass and NiO-doped bioactive glass which indicates that the addition of Ni2+ ion would not adversely affect the overall bioactivity and structure of the glass. The Ni2+ (0.58 Å) such as Co2+ (0.55 Å) and Mg2+ (0.57 Å) ions has similar ionic radii and

Fig. 11 Variation in shear modulus of glasses and glass-ceramics with NiO concentration


Fig. 12 Variation in bulk modulus of glasses and glass-ceramics with NiO concentration

cationic charge state, respectively. Because of the similar ionic radii and therefore similar charge-to-size ratio, it might be expected that Ni2+ would adopt similar structural role as Co2+ and Mg2+ ions and concluded that it acts as a network intermediate [21]. It was also mentioned herewith that due to charge/size ratio of Ni2+ ion, it would enter into [SiO4]4− network as [NiO4]2− species in tetrahedral coordination, but it could also act as network modifier in higher coordination. Azevedo et al. [43] had also made similar observation with Co2+ that due to same charge/ratio of Co2+ ion, it would enter into silicate [SiO4]4− network as [CoO4]2− complex ion in tetrahedral coordination but it would act

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as modifier in higher coordination. The role of nickel oxide in the bioactive glass and glass-ceramic would determine the connectivity (NC) and the bioactive glass properties such as ion release rates and its effect on hydroxyl carbonate apatite formation. The authors found that there was an excellent agreement between metal– oxygen (M–O) correlation obtained for Ni2+ by isotopic substitution and Ni/Co data sets by isomorphic substitution method which confirmed that both Ni2+ and Co2+ ion were isostructural in the bioactive glasses. The entry of heavier Ni2+ ion as [NiO4]2− tetrahedral species in place of lighter Si4+ ion in [SiO4]4− network has resulted the formation of Ni–O–Si bonds in silicate network structure, as such the mechanical properties and densities of the bioactive glass and glass-ceramics samples have increased considerably with increasing NiO concentration. Elastic properties of nickel oxide doped glass and glass-ceramic The values of longitudinal and shear ultrasonic wave velocities of base glass and nickel oxide substituted glass and its derivative are given in Table 3. The ultrasonic longitudinal and shear wave velocities increase with an increase in concentration of nickel oxide (0.41 to 1.65 mol%). The values were used to calculate the Young’s, shear and bulk modulus of elasticity, and Poisson’s ratio of glasses and glass-ceramics and are given in Table 4. For better representation, change in Young’s, shear and bulk moduli, and Poisson’s ratio with increasing the concentration of nickel oxide in the glasses and their ceramic derivatives are shown in Figs. 10, 11, 12, and 13, respectively [49]. The values for Poisson’s ratio were plotted within error bars against composition of nickel oxide substituted samples and given in Fig. 13. The results show that Young’s, shear, and bulk moduli increase with an increase in nickel oxide content in glass and glass-ceramics, whereas the Poisson’s ratio slightly decreases with increasing nickel oxide content. This increase in elastic modulus of bioactive glass and glass-ceramics is related to the increase in connectivity of the glass network, which resulted in an increase in the Young’s, shear, and bulk moduli respectively [24, 50, 51]. Scanning electron microscopic analysis of glass and glass-ceramic samples

Fig. 13 Variation in Poisson’s ratio of glasses and glass-ceramics with NiO concentration

The SEM micrographs of 45S5 glass and glass-ceramic as well as nickel oxide substituted glass and glass-ceramic samples before immersion in simulated body fluid (SBF) solution are shown in Figs. 14, 15, and 16. It is clear that the samples possess


J Aust Ceram Soc (2017) 53:939–951

Fig. 14 SEM micrographs for NiO glass samples before SBF treatment

irregular and asymmetrical grains of bioactive glass and glassceramic. Next, Figs. 15 and 17 represent the same samples after immersion in SBF for 15 days. It is visible that the samples were now enclosed by asymmetrical shape, and the grounded carbonated HA particles have grown into more than a few

agglomerates, consisting of spine-shaped hydroxyl carbonate apatite (HCA) layer. The micrographs show the formation of HCA on the surface of 45S5 and nickel oxide substituted bioactive glass and glass-ceramic samples after immersion in SBF solution for 15 days [52].

Fig. 15 SEM micrographs for NiO glass samples after SBF treatment for 15 days


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Fig. 16 SEM micrographs for NiO glass-ceramic samples before SBF treatment

Conclusions In the present investigation, a comparative investigation was made on measurement of the compressive, flexural strength, microhardness, and elastic properties of the nickel oxide substituted bioactive glass and glass-ceramic. Ni2+ ions were introduced in the glass composition for Si4+ ion in different

concentration for 0–1.65 mol% NiO to yield a chargebalanced (CB) series of bioglass melts. The Ni2+ ion might be present in the bioactive glasses in tetrahedral and octahedral coordinations, both depending upon the oxygen environment. The glass transition temperature (Tg) of bioglass samples were found to decrease with increasing NiO content in the glass. Moreover, all the glass-ceramic samples showed the

Fig. 17 SEM micrographs for NiO glass-ceramic samples after SBF treatment for 15 days


J Aust Ceram Soc (2017) 53:939–951

sodium calcium silicate, Na2Ca2 Si3O9, and Na2Ca2 Si3O8 as main crystalline phases. The SEM micrographs of the bioactive glass and glass-ceramic samples after SBF treatment for a definite time period have indicated the formation of HCA layer on the surface of the samples. This shows that the addition of NiO did not have any adverse effect on the overall bioactivity of the samples. Density, compressive strength, flexural strength, microhardness, and elastic moduli such as Young’s, shear, and bulk moduli were found to increase with increasing concentration of NiO in the glasses and their glassceramic derivatives. These destructive and non-destructive results show similar trends in behavior and are comparable in the systems, glass, and glass-ceramic. Hence, this investigation suggests that the mechanical behavior of the samples can be measured without any destruction of the sample, since the biomaterials are expensive to prepare. Acknowledgements The authors gratefully acknowledge the HOD Department of Ceramic Engineering, Indian Institute of Technology (Banaras Hindu University) Varanasi-221005, India and the honorable Director of Indian Institute of Technology (Banaras Hindu University) Varanasi, India for providing necessary facilities for the present work. The author, Vikash Kumar Vyas, is also very much grateful to the University Grants Commission, New Delhi, India (RGNF-SC-UTT-2012-13-25709) for providing the Rajiv Gandhi National Fellowship for the research work.

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