Chapter 2. 3D printing of shape-conformable thermoelectric materials using all-inorganic
2.2 Results and discussion
2.2.1 Rheological properties of all-inorganic TE inks
Figure 2.1a illustrates the process of extrusion-based 3D printing of TE materials of defined shapes using all-inorganic TE inks. As aforementioned, the major challenge for this strategy is to develop all- inorganic 3D printable TE inks with no organic binder, by tailoring rheology to ensure reliable flow through fine nozzles and structural integrity to withstand after deposition.16-19 Recently, inorganic ChaM ions have been reported to act as surface ligands of nano- and micro-scale particles, which stabilise the particles in solution via electrostatic interactions.30-35 In the present study, I found that the Sb2Te3 ChaM ions in the concentrated colloid inks effectively held Bi2Te3-based TE particles together in the electrostatic manner, which strongly affected their rheological properties and the related 3D printability (Figure 2.1b)
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Figure 2.1. (a) Photograph showing viscoelastic all-inorganic TE ink. (b) Illustration of an extrusion- based 3D printing process. (c) Optical microscopy image (top) and a photograph (bottom) of the 3D- printed TE materials. (d) Scheme of the structural changes of TE particles depending on the content of Sb2Te3 ChaM binders in the TE ink. Adding a small amount of the ChaM binders (yellow) induced a localized aggregation of BiSbTe particles (grey), reducing their effective volume. Further addition of the ChaM (orange) led to aggregates consisting of charged BiSbTe particles with a stronger electrostatic interaction. (e,f,) Dynamic viscosity, η′ (e) and loss tangent, tan δ (f) curves of the TE inks with the various ChaM contents (BST100/ChaM#, where # is the weight percentage to that of BiSbTe particles) at 25 °C. The dotted line indicates a tan δ value of 1 (liquid-like behaviour at tan δ > 1 or solid-like behaviour at tan δ < 1). (g) Variation of the yield stress (τy) and recoverable elastic strain (SR) as functions of the ChaM content in the BST100 ink. τy and SR represent the minimum energy to break down the physical structure and the extent of elastic recovery (the ratio of G′ (storage moduli) to 2G″ (loss moduli) at 0.05 rad s−1), respectively.SEM image of dried composite of BST and Ten2- polyanions. (B) XRD patterns of dried and annealed composite of BST and Ten2- at various temperatures.
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Figure 2.2. Time evolution of photographs showing the TE inks of (a) BST100/ChaM25 and (b) BST100/ChaM0.
The inks were prepared by the dispersion of Sb2Te3-based ChaM binders with desired amount in the suspension of Bi2Te2.7Se0.3 for n-type and Bi0.4Sb1.6Te3.0 for p-type TE particles in glycerol and the subsequent homogenization. The inks with different ChaM content are hereafter referred to as BST100/ChaM#, where ChaM# is the ChaM content expressed as the ratio of the weight of ChaM binders to that of BiSbTe particles in percentage. Figure 2.1c shows the dynamic viscosity (η') of the inks with a ChaM content of 0, 12.5, 25 and 37.5%. All the TE inks exhibited a noticeable Bingham behaviour, indicating the collapse of the colloidal structure under shear stress. At the low concentration of ChaM, the dynamic viscosity of the TE ink (BST100/12.5) is lower than that of BST100 over the whole frequency range. This suggests that adding a small amount of ChaM binders induced a localised aggregation of the charged BiSbTe particles, thus reducing their effective volume (Figure 2.1b). However, a further addition of ChaM (with a ChaM content higher than 25%) increased the η'. BST100/37.5 showed the highest η' of all four samples over the measured frequency
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range, suggesting that the aggregates consist of the charged BiSbTe particles experienced a stronger electrostatic interaction at a higher ChaM content (Figure 2.1b). The phase stability of the specimens was evaluated by loss tangent (tan δ) curves ((Figure 2.1d). An increase of tan δ indicates liquid-like character of the specimen, while a decrease indicates solid-like character.36 It is interesting to note that the tan δ value of the TE inks containing ChaM was below 1 in the lower frequency region and gradually increased with increasing frequency. Furthermore, a continuous decrease of tan δ was observed with increasing ChaM content, suggesting that the TE ink behaviour was more solid like and its colloidal stability was improved. On the other hand, BST100 (without any ChaM binder) yielded a tan δ value of 1.2 at 0.05 rad/s and showed shear-dependent fluctuation, which implies that the BST particles with no ChaM formed a liquid-like structure with poor colloidal stability. This is further supported by the fact that the ink with no ChaM frequently clogged the nozzle during printing. Further, BST100/25 displayed greater storage moduli (G') but lower tan δ values than BST100 (Figure 2.2).
These results demonstrate that the electrostatic binding effect of the ChaM improved the elasticity of the colloids, leading to a higher colloidal stability.
Figure 2.3. (a) The storage (G') and loss (G") moduli curves of the TE inks at 25 oC. (b) Variation of the storage modulus (G') and loss tangent (tan δ) values at 0.05 rad/s with the content of the ChaM in the TE ink.
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The strength and elastic behaviours of the TE inks can be quantitatively assessed by the yield stress (τy) and recoverable elastic strain (SR) as shown in Figure 2.1e, which represent the minimum energy to break down the physical structure and the extent of elastic recovery, respectively. The τy and SR values were obtained by the y-axis intercept of the Casson plot in Figure 2.3 and the ratio of G' to 2G" at 0.05 rad/s, respectively.37 In Figure 2.1e, the τy value of BST100 was 39.5 Pa, while that of BST100/25 was 49.4 Pa. On the other hand, the SR value of BST100/25 was 0.85, which is more than two times higher than that of BST100 (0.41), suggesting that the ChaM-containing TE inks would show better structural integrity upon printing than ChaM-free TE inks. This high elastic strain achieved in TE inks with the ChaM binders fulfil the desired viscoelasticity of inks for an extrusion- based 3D printing process. Regarding these rheological properties, printability, composition and related TE properties of the 3D printed TE materials, I chose the BST100/25 ink for my following 3D printing experiments.
2.2.2 3D printing of shape-conformable TE materials
Figure 2.4. Casson plots of the TE ink with the ChaM contents of 12.5, 25 and 37.5%, respectively, at 25 oC.
The n-type or p-type TE inks with the ChaM content of 25% were printed onto a carbon substrate through cylindrical nozzles with an inner diameter of 260 μm by an air-powered fluid dispenser (Figure 2.4). The viscoelastic properties of the inks enabled a continuous printing of the ink out of the
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nozzle under pressure. Also, the glycerol in the TE ink served as humectant and facilitated the integration of the individual printed layers into a single structure38, eventually enhancing its density, as revealed by the optical microscopy image shown in Figure 2.1.
Figure 2.5. A photograph showing the as-printed sample before drying from the TE ink with 10wt%
of the ChaM
Figure 2.6. Top-view SEM images of the 3D-printed n-type (a) and p-type materials (b). (c) XRD patterns of the 3D-printed n-type and p-type materials. The vertical lines indicate peaks corresponding to bulk Bi0.5Sb1.5Te3.0 (Joint Committee on Powder Diffraction Standards, JCPDS: 00-049-1713) and Bi2Te3 (JCPDS: 00-015-0863). (d) Photographs of the 3D-printed p-type cuboid, with widths ranging from 5 to 10 mm, disc and half-ring structures.
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After completing 3D printing process, the printed and dried n-type and p-type TE materials were heated at 450 oC under inert atmosphere for 1 h. This heat treatment process led to the sintering of TE materials with the Sb2Te3-based ChaM to obtain dense TE materials without the formation of secondary phases such as Sb2Te3 or Te. The scanning electron microscopy (SEM) images of the sintered samples (Figure 2.6a,b) reveals that the particles were well fused together into polycrystalline grains (Figure 2.5). The X-ray diffraction (XRD) patterns of the resulting n-type and p-type TE materials (Figure 2.6c) show the patterns corresponding to solely Bi2Te2.7Se0.3 and Bi0.5Sb1.5Te, thus excluding the formation of secondary phases such as Sb2Te3 or Te. The peaks were however slightly shifted to higher angles, which can be attributed to a slight change in the composition of the TE materials resulting from the integration of Sb2Te3-based ChaM into their host lattices during the heat treatment. At the same time, the sintering effect of the ChaM additive has been shown to cause a slight shrinkage of the 3D printed features.19,39,40 Typically, the dimensions of a sample were reduced by 20% in width and 20% in thickness from the initially printed 3D model. As this sintering shrinkage rate was highly uniform, it was possible to scale the CAD model accordingly before printing.
The major benefit of 3D printing is the shape engineerability and scalability of materials for preparing shape-conformable TE materials to heat sources. The currently developed 3D printing technology successfully demonstrates these capabilities by shaping a cuboid, a disc, and a half ring, as well as varying these objects’ width from 3.5 mm to 10.0 mm (Figure 2.6d). Furthermore, all these samples exhibit almost identical densities of 3.91 ± 0.10 g·cm-3. Since the density of 3D printed and sintered inorganic materials reflect the degree of shrinkage during these processes, these results further confirm the reproducibility of 3D printing process.
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Figure 2.7. The room-temperature electrical conductivity and Seebeck coefficient of the p-type 3D- printed disc, half ring and cuboid (a), cuboids with widths of 5.5, 8.0 and 10 mm (b) and three different pieces cut from the same half-ring (c). I characterized the room temperature electrical conductivities and Seebeck coefficients of more than three sets of the p-type 3D-printed samples and the uncertainties (standard errors) of the electrical conductivities and Seebeck coefficients were 0.8%
and 1.5%, respectively. Temperature-dependent TE properties of the n-type and p-type 3D-printed cuboids: electrical conductivity and absolute Seebeck coefficient (d), thermal conductivity (e) and zT (f). The stars indicate data points obtained from the hot-pressed Bi2.0Te2.7Se0.3 (n-type)31 and Bi0.4Sb1.6Te3.0 (p-type)45 materials. A photograph showing the as-printed sample before drying from the TE ink with 10wt% of the ChaM
Figure 2.8. A photograph of the extrusion-based 3D printing equipment by an air-powered fluid dispenser.
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Figure 2.9. SEM image of the n-type 3D-printed TE materials upon drying at 105 oC. Scale bar corresponds to 10µm.
The 3D printing processability and reproducibility in the current technology are clearly reflected by the reproducible TE properties of the printed samples with various shapes and dimensions. The room temperature electrical conductivities and Seebeck coefficients of the p-type 3D printed cuboid, disc, and half ring are 55375 S/m and 165 μV/K with the uncertainty of 0.8~1.5% (Figure 2.7a).
These high electrical properties are comparable to those of BiSbTe materials obtained by the conventional procedures.12 Furthermore, the cuboids with a width ranging from 5.0 mm to 10.0 mm exhibited the almost identical electrical conductivity and Seebeck coefficient (Figure 2.7b). Another prominent issue in shaping TE materials into complicated structures is the inhomogeneity of TE properties.41 To demonstrate the TE property homogeneity in 3D printed TE materials, the TE half ring was cut into three pieces, and the electrical conductivity and Seebeck coefficient (Figure 2.7c) were found to be the same for all three samples, and no degradation in the electrical properties of the cut samples was observed in comparison to the original half ring, demonstrating the homogeneity of the TE properties.
The temperature dependence of the TE properties was characterised on the n-type and p-type 3D printed cuboids with a width of 10 mm and a thickness of 1.5~2.0 mm. The electrical conductivity of the n-type and p-type samples was in the range 50000–55000 S·m-1, and it decreased with increasing temperature (Figure 2.7d). The Seebeck coefficient of the n-type and p-type samples increased with increasing temperature, with a peak value of 145 and 199 μV·K-1 for the n-type and p-type samples at 175~200 oC, respectively (Figure 2.7d). The carrier mobility and concentration of these samples were determined at room temperature by a Hall effect measurement. The n-type and p-type samples exhibited a carrier mobility of 78.5 and 93.6 cm2·V-1·s-1, respectively. These high mobility values account for the high electrical conductivity of the 3D printed samples. On the other hand, the carrier
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concentrations of the n-type and p-type samples were 3.99 × 1019 and 3.64 × 1019 cm-3, respectively.
These values are relatively high compared with those of typical bulk Bi2Te3-based materials. Since the Seebeck coefficient is reciprocally proportional to the carrier concentration, the slightly low Seebeck coefficient of the 3D printed TE samples compared with the corresponding bulk values is attributed to the high carrier concentration.3
The temperature-dependent thermal conductivity of the n-type and p-type printed samples was 0.50–0.63 W·m-1·K-1 over the entire measured temperature range (Figure 2.7e), which was significantly reduced in comparison to 1.5–2.5 W·m-1·K-1 of bulk Bi2Te3-based materials.12 In addition, the thermal conductivity showed a weak negative temperature dependence (Figure 2.7e), which can be attributed to a weak bipolar contribution to the thermal conductivity due to the relatively high carrier concentrations. To further understand the thermal properties of the 3D printed samples, the lattice thermal conductivity (κL) of these samples was calculated by subtracting the electronic contribution of thermal conductivity (κE) from the total thermal conductivity, where κE was calculated by the Wiedemann-Franz Law, κE=L0σT, where L0 is the Lorenz number, σ is the electrical conductivity and T is the absolute temperature. Also, L0 was estimated by the equation based on the single parabolic band (SPB) model. The calculated values of κL of n-type and p-type samples ranges from 0.28 W m-1 K-1 to 0.36 W m-1 K-1 (Figure 2.11) which are slightly lower than the theoretically predicted lowest limit of the lattice thermal conductivity of Bi2Te3-based materials (0.21 W m-1 K-1).
These ultralow thermal properties can be understood by considering the porosity of the 3D printed materials. I reported the paintable TE materials in that the painted and sintered TE layers exhibited the even lower lattice thermal conductivity that the currently 3D printed materials35. The painted materials had the multi-scale pores in nano- and macro-scales, which effectively scatter phonons in a broad range of wavelength to reduce the lattice thermal conductivity. Meanwhile, the macro-scale pores less affected the electrical carrier transport due to the relatively short mean free path of electron and hole carriers. Likewise, the current 3D printed and sintered TE materials also possess the macro-scale pores to scatter phonons (Figure 2.12), responsible for the low lattice thermal conductivity. Also, the relatively high carrier mobility of the 3D printed TE materials further supported the existence of macro-scale pores.
To quantitatively understand the porosity effect on the κL, the κL of the n-type and p-type 3D printed samples were applied to the modified formulation of the effective medium theory equation of 𝜿𝑳 = 𝜿𝒉(𝟐−𝟐𝚽)
(𝟐+𝚽) suggested by Lee, et al., where κh and Φ are the lattice thermal conductivity of host materials and the porosity, respectively, The estimated κL were 0.62 W m-1 K-1 for the n-type sample
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and 0.68 W m-1 K-1 for the p-type sample (Figure 2.7d), which were in the similar range of those observed in the nanostructured thermoelectric Bi2Te3 and BiSbTe materials (Figure 2.10)27.
These promising electrical and thermal properties lead to the remarkably high zT values of n-type and p-type 3D printed samples. The n-type and p-type 3D printed samples exhibited a zT value of 0.4 and 0.7 at room temperature, respectively (Figure 2.7f). Furthermore, the maximum zT values of 0.6 and 0.9 were achieved at 170 oC and 125 oC for n-type and p-type TE materials, respectively. These values are comparable to those obtained for typical Bi2Te3 and BiSbTe ingots (zT ~0.8-1.0)3, and much higher to the value of 0.12 reported for amorphous TE materials 3D printed by
stereolithography (SLA).43 To the best of my knowledge, the zT values reported in this study are among the highest for TE materials produced from TE inks or pastes.35,42,44
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Figure 2.10. Comparison of TE properties of the 3D printed samples with the hot pressed ones with corresponding compositions and (c) thermal conductivity. The stars in panel (c(a) Electrical conductivity, (b) indicate data points obtained from the hot Bi2.0Te2.7Se0.3 and Bi0.4Sb1.6Te3.0.) Seebeck coefficient, pressed.
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Figure 2.11. Lattice thermal conductivities of and corrected lattice thermal conductivities effective medium theory the n-type and p-type 3D printed using the modified formulation of the (dotted lines).
Figure 2.12. SEM image of fractured surfaces of (a) n-type and (b) p-type 3D-printed TE materials annealed at 450oC. Scale bar corresponds to 10 μm.
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2.2.4 Conformal cylindrical TEGs with 3D printed TE half rings
Figure 2.13. a, Scheme showing the fabrication of the TEG. Cu electrodes (brown) were attached on the top of an alumina pipe with inner and outer diameters of 5 mm and 8 mm. Then, three pairs of n- type (green) and p-type (orange) TE half-rings materials (8 mm inner diameter, 15 mm outer diameter,
~1.53 mm average thickness) were 3D-printed and attached to Cu electrodes using Ag-containing epoxy. Finally, top Cu electrodes were assembled with Ag-containing epoxy. b, Scheme for the measurement of output power under flowing hot water through an alumina pipe. For measuring the TEG output characteristics, hot water was constantly circulated inside the alumina pipe using an electric water pump. c, Illustration of the TEG for the measurement. Two T-type thermocouples and two electrical leads were connected to the TEG to measure the temperature difference and the output voltage. d, Photographs of the fabricated half-ring-based conformal TEG. The inset shows the 3D- printed n-type and p-type half rings. e, Output voltage and power of the cylindrical TEGs at different temperature differences.
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Figure 2.14. Three different pieces cut from the same half-ring for measuring TE properties.
Figure 2.15. Schematic showing the temperature distributions of the half ring-based conformal TEG (a) and the conventional planar TEG calculated by the FEM (b). (c) Output voltage per thermocouple (top) and output power (bottom) of the conformal cylindrical TEG as a function of the temperature difference obtained by FEM (line) and experiment (symbol). Calculated output voltage per thermocouple (top) and output power (bottom) (d) and calculated heat rates dissipated from a water tube and absorbed by the TE legs (top) and generating efficiency (bottom) (e) of the half ring-based conformal TEG and the conventional planar TEG as a function of the temperature difference.
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The remarkable TE properties and 3D printability enable to design of TE materials conformable to the shape of heat sources, which can be the most effective means to transfer heat energy from a heat source to a TEG. In this study, I propose a fully-integrated conformal cylindrical TEGs to a heat source, made of 3D printed TE half rings.4,8 This conformal TEGs were prepared by using 3D printed TE half rings and mounting them on an alumina pipe. Figure 2.13a illustrates the overall fabrication process of the TEG composed of three pairs of n-type and p-type half-ring TE legs with an inner diameter of 8 mm that perfectly fitted the alumina pipe (Figure 2.13b). The TE legs were attached on thin Cu electrodes mounted on the alumina pipe by using Ag epoxy. The overall module resistance (RTEG) was 105 mΩ, i.e. ~5 times higher than the calculated resistance of 20 mΩ. This difference is attributed to the electrical contact resistance (ρc) between the 3D printed TE legs and Ag inter-layers since the contact resistance between Bi2Te3-based materials and Ag electrodes fabricated from Ag paste or paint was reported to be an order of magnitude higher than that for modules bonded by soldering.35 However, the conventional process of soldering presents a technical challenge in the present set-up since the 3D printed half-ring-based TE legs would make the melted solder flow down their curved surface. This is the reason why I used an electrically conductive Ag epoxy paste to connect the TE half-rings to the Cu electrodes. Despite its higher contact resistance, this paste provided an excellent processability and mechanically robust contacts between the electrodes and TE legs.
The output characteristics of the conformal cylindrical TEGs were measured under a flow of hot water through the alumina pipe. By controlling the hot water temperature, the temperature at the hot side of the surface of the alumina pipe was well modulated, while the cold-side temperature was maintained at 25~30 oC (Figure 2.14). As the temperature difference increased, the TEGs showed a linear and quadratic increase of the output voltage and power, respectively, achieving a maximum output voltage of 27.0 mV and maximum power of 1.62 mW at a temperature difference of 39 oC (Figure 2.13c).
To show the effectiveness of the half-ring-based conformal TEG, I developed a three-dimensional finite element model (FEM) that allowed us to calculate the temperature and electrical potential distributions in the TEGs under a flow of hot water through the alumina pipe (Figure 2.15a,b).
Furthermore, I compared the generation from the current conformal TEG and the conventional planar TEG mounted on the alumina pipe under the same environmental condition. For this comparison, a planar type of TEG with a similar substrate area to the half-ring-based conformal TEG is considered.
The temperature gradient within the TE elements (ΔTTE) was larger with the half-ring-based TEG (ΔTTE ~ 29 K) than with the planar TEGs. The planar TEG with an epoxy width (wEpoxy) of 4 mm exhibited a larger ΔTTE (9 K) than the one with wEpoxy of 1 mm (8 K).