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Interpreting the combustion process for high performance ZrNiSn thermoelectric materials Tiezheng Hu, Dongwang Yang, Xianli Su, Yonggao Yan, Yonghui You, Wei Liu, Ctirad Uher, and Xinfeng Tang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b15273 • Publication Date (Web): 13 Dec 2017 Downloaded from http://pubs.acs.org on December 15, 2017

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ACS Applied Materials & Interfaces

Interpreting the combustion process for high performance ZrNiSn thermoelectric materials Tiezheng Hua, Dongwang Yanga, Xianli Sua*, Yonggao Yana, Yonghui Youa, Wei Liua, Ctirad Uherb and Xinfeng Tanga* a

State Key Laboratory of Advanced Technology for Materials Synthesis and

Processing, Wuhan University of Technology, Wuhan 430070, China b

Department of Physics, University of Michigan, Ann Arbor, MI 48109, USA

KEYWORDS: thermoelectric, half-Heusler, ZrNiSn, SHS, thermoelectric properties

ABSTRACT: ZrNiSn alloy, a member of the half-Heusler family of thermoelectric materials, shows great potential for mid-to-high temperature power generation applications due to its excellent thermoelectric properties, robust mechanical properties, and good thermal stability. The existing synthesis processes of half-Heusler alloys are, however, rather time and energy intensive. In this study, single-phase

ZrNiSn

bulk

materials

were

prepared

by

self-propagating

high-temperature synthesis (SHS) combined with spark plasma sintering (SPS) for the first time. The analysis of thermodynamic and kinetic processes shows that the SHS reaction in the ternary ZrNiSn alloy is different from the more usual binary systems. It consists of a series of SHS reactions and mass transfers triggered by the SHS fusion of the binary Ni-Sn system that eventually culminates in the formation of single-phase 1

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ternary ZrNiSn in a very short time, which reduced the synthesis period from few days to less than an hour. Moreover, the non-equilibrium feature induces Ni interstitials in the structure, which simultaneously enhances the Seebeck coefficient and decreases the thermal conductivity, favorable for thermoelectric. The maximum thermoelectric figure of merit ZT of the SHS+SPS-processed ZrNiSn1-xSbx alloy reached 0.7 at 870 K. This study opens a new avenue for the fast and low-cost fabrication of half-Heusler thermoelectric materials.

1. Introduction Thermoelectric (TE) materials can accomplish a direct reversible conversion between thermal energy and electricity via the Seebeck and Peltier effects. TE energy conversion has attracted great attention in the areas of industrial waste heat recovery and spot cooling of electronic devices. Its chief attributes are exceptionally reliable, noiseless, and environmentally friendly operation. The conversion efficiency of thermoelectric materials is determined by their dimensionless figure of merit ZT defined as ZT = α2σT/κ, where α, σ , κ and T are the Seebeck coefficient, the electrical conductivity, the thermal conductivity and the absolute temperature, respectively1-2. Half-Heusler (HH) alloys have drawn a lot of attention due to their nontoxic elements and excellent thermoelectric and mechanical properties 3-6. Among them, the ZrNiSn-based HH alloy possesses excellent electronic transport properties due to its 18-valence electron structure and sharp density of states near the Fermi level7-10. Very good mechanical properties and thermal stability are also highly 2


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desirable for high temperature power generation mainly synthesized either by arc melting (AM) levitating melting (LM)

9-12

. So far, HH alloys have been

13-14

, by induction melting

17-18

, or by long time solid state reactions

15-16

, by

19-20

. However, the

equipment for arc melting, induction melting and levitating melting are relative expensive, and solid-state reaction is a time and energy consuming process. Therefore, it is crucial to develop a facile and cost-efficient fabrication method for HH materials which can shorten the synthesis period with minimum energy consuming. Self–propagation high–temperature synthesis (SHS), which is also referred to as the combustion synthesis (CS), is a novel process for rapid preparation of materials by using the energy released during the chemical reaction. The SHS synthesis process can be finished in a few seconds with consuming minimum energy once the reaction is initiated

21-24

. Moreover, the SHS process is so fast that the evaporation of low

melting point elements is dramatically suppressed, resulting in a precisely controlled composition. In the past, the SHS technique has been used exclusively to synthesize high temperature refractory materials, such as ceramics, composites and intermetallic compounds that required achieving the adiabatic temperature in excess of 1800 K to make the reaction self-sustaining, the so-called Merzhanov’s criterion

23

. Because

most of the thermoelectric materials would either melt or decompose at the temperature above 1800 K, essentially no attempts were made to apply the SHS technique to their synthesis. Recently, we have successfully applied the SHS synthesis to many thermoelectric compounds, including Cu2Se 25, Bi2Te3 26-27, 31, skutterudites 28, BiCuSeO 29, and SnTe 30. It not only dramatically shortened the fabrication time, but 3



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also generated numerous non-equilibrium microstructure (nanostructure, interstitial or meta-stable phase) in the products due to non-equilibrium feature of SHS processing, which would impede the thermal transport, lowering the thermal conductivity. Thus, the structure also possessed excellent thermoelectric properties25-30. Our investigations have, so far, focused on the synthesis of binary thermoelectric materials. With the expertise and understanding gained, we have decided to tackle a more challenging task: attempting to use the SHS synthesis in the preparation of ternary thermoelectric compounds with non-equilibrium structure. In this work, we apply the SHS process to a preparation of ZrNiSn, an important member of the half-Heusler family of thermoelectric materials. The SHS process of ternary ZrNiSn includes a series of SHS reactions and mass transfers initiated by fusion of Ni and Sn, and culminating in a synthesis of a single-phase ternary ZrNiSn compound in a very short time. We describe the relevant thermodynamic and kinetic processes, phase transformations, and we correlate them with the microstructural evolution. Due to the rapid, non-equilibrium nature of the SHS process, it induces Ni interstitials in the structure which increase the electrical conductivity and suppresses the thermal conductivity effectively. The figure of merit ZT of the ZrNiSn alloys prepared by SHS-SPS reached 0.67 at 870 K, which is comparable with the figure of merit of ZrNiSn prepared by the traditional methods10,

14

. The maximum

thermoelectric figure of merit ZT of the SHS+SPS-processed ZrNiSn1-xSbx alloy reached 0.7 at 870 K.

4


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2. Experiment Commercial high-purity powders of Zr (3.5 N, 200 mesh), Ni (2.5 N, 200 mesh), Sn (2.5 N, 200 mesh) and Sb (3.5 N, 200 mesh) were weighed according to a stoichiometric ratio of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03), and then mixed uniformly in an agate mortar. The mixtures were cold pressed into pellets with the diameter of about 15 mm. A small hole with a diameter of 1 mm was drilled to a depth of about 5 mm at the bottom of the pellet to accommodate a thermocouple, which was used to measure the combustion temperature. Monitoring and recording of the SHS process was carried out in a home-built equipment shown in Figure 1(a). Specifically, the combustion temperature and the propagation velocity of the combustion wave were conveniently measured in this apparatus. For the ZrNiSn alloy, the SHS reaction proceeded under the protective atmosphere of 50 kPa pressure of argon. The intent was to further suppress volatilization of Sn and Sb. The pellet was ignited by a beam from a continuous wave fiber laser having the diameter of 100 µm and the wavelength of 1064 nm. The ignition powers used were 50 W and 150 W, respectively. An additional thermocouple was inserted into the center of the pellet to monitor the combustion temperature evolution. Thermocouple readings were collected by a digital multimeter (Agilent 34420A). The reaction process was also observed using a high-speed camera (OLYMPUS i-SPEED 3) with the frames advancing at 400 fps. The propagation velocity of the combustion wave was processed by software. The products obtained after SHS were ground into fine powders and then SPSsintered under a pressure of 30 MPa at 1273 K with the holding time of 5 min. ZrNiSn1-xSbx (x = 0, 0.01, 0.02, 0.03) bulk materials with a relative density above 98% were achieved. The sintered pellets were cut into sizes suitable for transport measurements. Following the initial transport evaluation, the samples were annealed 5



at 1073 K for 7 days and their transport properties were re-measured to evaluate their thermal stability. The structure of SHS-prepared powders and the SPS-sintered bulks were examined by a powder X-ray diffractometer (PANalytical: Empyrean, Cu Kα). The morphology and elemental distribution were determined from back-scattering electron images (FESEM, SU8020) and energy dispersive X-ray (EDX) analysis (Bruker), respectively. The actual compositions of un-annealed ZrNiSn1-xSbx are determined by the wavelength dispersive spectroscopy (WDS) Electron probe micro-analyzer (EPMA, JXA-8230). The electrical conductivity and the Seebeck coefficient were measured simultaneously by a standard four-probe method (ULVAC-RIKO ZEM-3) in a helium atmosphere. The thermal conductivity was calculated from the measured thermal diffusivity D, specific heat Cp, and density ρ according to the relationship κ = DCpρ. The thermal diffusivity and the specific heat were determined by a laser flash method using a NETZSCH: LFA 457 apparatus and a power compensation differential scanning calorimeter TA: DSC Q20 in an argon atmosphere, respectively. All measurements were performed in the temperature range from 300 K to 900 K. The actual sample density was determined by the method of Archimedes. The low temperature Hall coefficient measurements (10 - 300 K) were carried out in a Physical Properties Measurement System (PPMS-9, Quantum Design), making use of a five-probe sample configuration while sweeping the magnetic field between -1.0 T and 1.0 T. The carrier concentration (nH) and the room temperature Hall mobility (µH) were obtained from the Hall coefficient (RH) and the electrical conductivity by the relation: n =1/e|RH| and µH = σ|RH|, respectively, where e is the elemental electron charge.

6


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3. Result and Discussion Using equipment described in experimental section, and referring specifically to Figure 1(a), the high-speed camera recorded the SHS reaction taking place in ZrNiSn. Figures 1(b) and 1(c) show different stages of the SHS process ignited at the top of the pellet with different laser powers. Figure 1 (b) depicts the SHS process in ZrNiSn ignited with the laser power of 50 W. Two combustion wave fronts are observed. The initial combustion wave is rather weak. About 2 seconds right after the first combustion reaction ignited, the second combustion reaction starts accompanied with intense heat released and strong light emission, which is much stronger than the first combustion reaction. The combustion wave propagated through the whole pellet very quickly. Due to the massive amount of heat released during the SHS reaction, the pellet eventually collapsed. After analyzing the SHS process video, the respective propagation velocities of the two combustion wave fronts were 2 mm/s and 50 mm/s, within the range of reported values of 1–150 mm/s 31, indicating that both combustion wave fronts represent a typical SHS process. A question arises: why two combustion processes were observed? We believe this is related to the low laser power of 50 W used to ignite the reaction. Such low laser power merely initiated the reaction with low ignition temperature. As the reaction slowly proceeded, the pellet’s temperature increased rapidly and, once it reached the minimum temperature necessary to trigger the second combustion reaction, a much more vigorous reaction took place. Since the propagating velocity of the second combustion reaction is much larger than that of the first one, once the second SHS reaction was ignited, the combustion wave of the 7



second SHS reaction caught up with the first SHS reaction, and the combustion wave fronts passed through the entire pellet simultaneously. To verify our hypothesis, we subsequently employed higher laser power of 150 W to ignite the SHS process. Figure 1(c) shows different stages of the SHS process in ZrNiSn ignited with the laser power of 150 W. Indeed, with this much higher laser power of 150 W, only one propagating combustion wave front is detected moving with the speed of 50 mm/s, and the reaction is very intense. Obviously, with this larger laser power of 150 W, both combustion reactions are ignited virtually simultaneously and only one (combined) combustion wave front is observed. A thermocouple inside the ZrNiSn pellet recorded the temperature rise shown in Figure 1(d). Due to a strongly exothermic nature of the combustion process, the combustion temperature rises very quickly, within a few seconds, to 1770 K and then slowly cools down to room temperature. The XRD pattern of the product after SHS is shown in Figure 1(e). A perfect match of the XRD pattern with the standard PDF card of ICSD# 98-064-6828 attests to the formation of a single-phase ZrNiSn half-Heusler alloy with no visible trace of any impurities during a very short time duration of the SHS process.

8


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Figure 1. (a) Schematic diagram of an apparatus to observe the SHS process; (b) Different stages of the SHS process in ZrNiSn when the laser ignition power was 50 W; (c) Different stages of the SHS process in ZrNiSn when the laser ignition power was 150 W; (d) Temperature profile of the SHS process in ZrNiSn; (e) XRD patterns of the product after SHS and after SHS-SPS.

To glean deeper into the phase transformation mechanism and microstructural evolution taking place in ZnNiSn during the SHS process, we used a technique of combustion wave front quenching method which relies on a rapid quenching (extinguishing) of the combustion wave front to retain the characteristics of the reaction at that particular stage and in different sections of the pellet. The technique

9



was first used by Rogachev et al to study phase transformations taking place during the SHS process of TiC and TiB2

32

. By analyzing XRD and FESEM images of

different parts of the quenched product, correlations between phase transformations and microstructural evolution at different stages of the combustion process can be evaluated. To apply the quenching technique, we prepared new cold-pressed ZrNiSn pellets as described before, except that the pellets were pressed in a smaller steel die of 10 mm diameter. Due to the very intense SHS reaction in ZrNiSn, the steel die was soaked in liquid nitrogen to conduct the heat away faster, otherwise the progression of the combustion wave front could not be interrupted. A schematic diagram of the set up to quench the combustion wave front is shown in Figure 2(a). Figure 2(b) is a schematic diagram of the quenched product divided into an unreacted area, a reaction area and a product area. To make it clearer, the reaction area is further subdivided into reaction area I, reaction area II and reaction area III.

Figure 2. (a) Schematic diagram of the combustion front quenching device; (b) 10


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Schematic diagram of the product following combustion wave front quenching.

Table 1 summarizes the XRD results for different sections of the quenched sample. XRD results show that, the unreacted area consists of mainly elemental Zr, Ni and Sn. In reaction area I, elemental Zr, Ni and Sn are still detected, but Ni3Sn2 and Ni3Sn4 phases begin to show up. In reaction area II, elemental Ni and Sn have totally disappeared, forming Ni3Sn2 and Ni3Sn4 phases, and the ZrNiSn phase begins to show up. In reaction area III, a new ZrNi2Sn full-Heusler phase together with elemental Sn are detected. In addition, the Ni3Sn4 phase has disappeared and the presence of the Ni3Sn2 phase has dramatically diminished. Finally, in the product area, a single phase ZrNiSn is obtained.

Table 1: Summary of the XRD results for different areas of the quenched sample.

Area

XRD Results

Unreacted area

Zr, Ni, Sn

Reaction area I

Zr, Ni, Sn, Ni3Sn2, Ni3Sn4

Reaction area II

Zr, ZrNiSn, Ni3Sn2, Ni3Sn4

Reaction area III

ZrNiSn, ZrNi2Sn, Zr, Sn, Ni3Sn2

Product area

ZrNiSn

We have already pointed out that with the low laser ignition power of 50 W, the combustion process consisted of two reactions: a relatively mild first reaction propagating with 2 mm/s and a much more intense second reaction propagating with 50 mm/s. To clarify the relationship between the two combustion processes, how the intermediate phases form, and how the phase transformations eventually lead to the 11



final single-phase ZrNiSn, we studied binary combustion reactions with molar ratio of 1:1 among elemental Zr, Ni and Sn. The results show that reactions between Zr and Ni and between Zr and Sn are very intense with a lot of heat released. In contrast, the reaction between Ni and Sn is relatively mild. Thus, the first step in the ZrNiSn combustion process is this mild combustion reaction between Ni and Sn, consistent with the results of combustion wave front quenching. The reaction between Ni and Sn is a classical SHS reaction with the propagating velocity 2 mm/s and the combustion temperature of 1000 K, exactly what we observed with the ignition laser power of 50 W. Subsequently, we mixed the product of the binary reaction between Ni and Sn with elemental Zr in the ratio Zr:Ni:Sn=1:1:1, and carried out a combustion reaction. The XRD pattern of the reacted products is shown in Figure 3(a). The product of binary Ni and Sn reacted in the ratio of 1:1 is a mixture of Ni3Sn4 and Ni3Sn2 compounds. The combustion reaction between Zr and the reacted products of Ni and Sn is also a classical SHS reaction, proceeding with the propagating velocity of 40 mm/s and attaining the combustion temperature of 1673 K, both parameters being smaller than the parameters characterizing the combustion reaction starting with elemental Zr, Ni, and Sn and ignited with the laser power of 50 W. SHS reactions between Zr and Ni3Sn2, Ni3Sn4 were carried out to find out how the binary Ni and Sn products react with Zr. XRD patterns of the products after SHS are shown in Figure 3(b). The products of Zr and Ni3Sn4 after the SHS reaction are mixtures of ZrNiSn with elemental Sn, while the products of Zr and Ni3Sn2 after the 12


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SHS reaction are ZrNiSn and ZrNi2Sn. Therefore, in reaction area III, ZrNiSn, ZrNi2Sn, and elemental Sn are detected. To reveal the phase transformation process between ZrNi2Sn, Zr and Sn, we ignited an SHS reaction between Zr, Sn and ZrNi2Sn. Following this SHS reaction, a single phase ZrNiSn half-Heusler alloy was obtained, as shown in Figure 3(b). According to the above analysis, the phase transformation mechanism during the SHS reaction of ZrNiSn is characterized by the following set of reactions: Ni + Sn → Ni3Sn4 + Ni3Sn2

（1）(SHS)

Zr + Ni3Sn4 → ZrNiSn + Sn

（2）(SHS)

Zr + Ni3Sn2 → ZrNi2Sn + ZrNiSn

（3）(SHS)

Zr + Sn + ZrNi2Sn → ZrNiSn

（4）(SHS)

Figure 3. (a) XRD patterns of products of binary Ni and Sn after combustion reaction and XRD patterns of the products of Ni and Sn after reacting with Zr via SHS; (b) XRD patterns of the products of Zr reacted with Ni3Sn2, Ni3Sn4 and ZrNi2Sn, respectively, by SHS. 13



To explore the microstructural evolution, we applied FESEM and EDS analyses on different areas of the quenched pellets, see Figure 4. Figure 4(a) depicts the microstructure of the unreacted area; Figs. 4(b) and 4(c) show reaction area I; Figure 4(d) corresponds to reaction area II; Figure 4(e) shows reaction area III; and Figure 4(f) shows the microstructure of the product area. The combined results of FESEM and EDS show that there are four stages of the microstructure evolution. In the unreacted area in Figure 4(a), one observes raw elemental Zr, Ni and Sn particles with the grain size ranging from a few microns to tens of microns. During the initial stage of the SHS process, corresponding to the reaction area I, Sn melts and coats the surface of the other particles under capillary force. Ni rapidly dissolves in molten Sn, and it is difficult to find elemental Ni. Figure 4(b) depicts numerous molten Sn-rich phases, consistent with the EDS result. In addition, isolated elemental Zr particles are confirmed by EDS. During the intermediate stage of the SHS process, corresponding to the reaction area II, Sn reacts with Ni, forming Ni3Sn4 and Ni3Sn2. As shown in Figure 4(c), many massive structures are attached to the surface of Zr. EDS results show that the massive structures are Ni-Sn phases (the mixture of Ni3Sn4 and Ni3Sn2). Subsequently, Ni3Sn4 reacts with Zr to form ZrNiSn and Sn, and Ni3Sn2 reacts with Zr to form ZrNiSn and ZrNi2Sn. As shown in Figure 4(d), the Ni-Sn phase, the ZrNiSn phase and the ZrNi2Sn phase are mixed together. Due to their similar microstructures, they can be distinguished only through the EDS analysis. During the final stage of the SHS process, corresponding to the reaction area III, ZrNi2Sn reacted with Zr and Sn to form ZrNiSn. Figure 4(e) shows the ZrNiSn phase and the ZrNi2Sn phase as mixed 14


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together. The area where the entire phase is ZrNiSn is characterized by the inter-granular fracture, while the area where the ZrNi2Sn phase dominates is characterized by the trans-granular fracture. Finally, after all reactions have been completed, single-phase ZrNiSn is obtained, as shown in Figure 4(f). EDS confirms that the entire area corresponds to the single-phase ZrNiSn half-Heusler alloy.

Figure 4. Microstructure of different areas of the quenched sample (a) Unreacted area; (b-c) Reaction area I; (d) Reaction area II; (e) Reaction area III; (f) Product area

Figure 5(a) shows a FESEM image of the fractured surface of ZrNiSn after SHS. The grain size is in the range of 2-10 µm, and the fracture can be characterized as 15



inter-granular. Figure 5(b) shows a FESEM image of the fractured surface of ZrNiSn after SPS. The sample is fully condensed and the fracture mode is again inter-granular. Figure 5(c) shows a BSE image of the polished surface of bulk ZrNiSn after SHS-SPS. The bulk sample is homogeneous with no trace of any secondary phase. Elemental distribution maps in Figs. 5(d)-5(f) show all elements distributed homogeneously. We conclude that single-phase and fully dense bulk samples are achieved via the SHS-SPS process in a very short time of less than one hour.

16


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Figure 5. (a) FESEM image of the ZrNiSn product after SHS; (b) FESEM image of ZrNiSn bulk after SHS-SPS; (c) BSE image of ZrNiSn after SHS-SPS; (d-f) Elemental distribution map of ZrNiSn bulk after SHS-SPS.

Table 2 Actual compositions, Lattice parameter, carrier concentration nH, carrier mobility µH, Seebeck coefficient α and lattice thermal conductivity κL of ZrNiSn1-xSbx at room temperature. ZrNiSn1-xSbx

Actual composition

Lattice parameter

nH

µH

1020 cm-3 cm2V-1s-1

α

κL

µVK-1

Wm-1K-1

x=0

ZrNi1.02Sn1.01

6.1142 Å

0.6

21.2

-248

5.9

x=0.01

ZrNi1.01Sn0.99Sb0.009

6.1134 Å

3.4

28.3

-130

6.2

x=0.02

ZrNi1.03Sn0.97Sb0.02

6.1131 Å

6.1

28.3

-97

6.8

x=0.03

ZrNi0.99Sn0.95Sb0.029

6.1123 Å

7.9

30.7

-79

6.7

Table 2 summarized the actual compositions, Lattice parameter, carrier concentration nH, carrier mobility µH, Seebeck coefficient α and lattice thermal conductivity κL of ZrNiSn1-xSbx at room temperature without annealing. The WDS composition analysis shows that SHS process can precisely control the composition. The calculated the lattice parameter of Sb doped ZrNiSn1-xSbx without annealing by Rietveld method is almost unchanged as the Sb content increase, due to the similar size of Sn and Sb. Figure 6(a) shows the temperature dependent carrier concentration measured in ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys prepared by the SHS-SPS method. The carrier concentration of un-annealed ZrNiSn first increases with the increasing 17



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temperature approaching the maximum carrier concentration of 3.33×1021 cm-3 at about 50 K, and then the carrier concentration decreases with the increasing temperature. At temperatures above 250 K, the carrier concentration becomes almost constant and independent of temperature. The room temperature carrier concentration of SHS-SPS prepared ZrNiSn is 5.8×1019 cm-3, comparable to that of the LM-SPS prepared sample with 5.0×1019 cm-3 10 and higher than that of the AM-SPS prepared sample having the carrier concentration of 3.0×1019 cm-3

14

. This abnormal

temperature dependent carrier concentration indicates possible presence of impurity states in the un-annealed ZrNiSn sample. The carrier concentration of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) is very sensitive to the presence of Ni interstitials33 in the structure and the Sb dopant content 10. Previous experimental and theoretical studies show that Ni can form Ni interstitial in the structure forming the impurity states which would increase the carrier concentration and lead to a distortion of electronic density states

34

. Because the non-equilibrium nature of the SHS process (high heating rate,

high cooling rate), Ni interstitials might form after SHS, which would induce the impurity states, leading to an increase of the carrier concentration with increase of temperature at the temperature below 50 K due to the ionization of the impurity states. Moreover, with the increasing level of Sb doping, the carrier concentration increases due to the out-shell electron difference between Sb and Sn. The carrier concentrations of annealed ZrNiSn, un-annealed ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) and annealed ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) are roughly independent of temperature in the measured temperature range and the carrier concentrations of all un-annealed samples 18


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are higher than those of the annealed samples with the same nominal composition， behaving as a highly degenerate semiconductor, revealing that doping with Sb would inhibit the formation of Ni interstitials.

After annealing, the concentration of Ni

interstitials or defects in the lattice decreases very dramatically and the carrier concentration correspondingly decreases. Figure 6(b) shows the temperature dependent carrier mobility of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys prepared by the SHS-SPS method. For un-annealed ZrNiSn, the carrier mobility at first decreases with the increasing temperature and then increases as the temperature climbs above 50 K. At 300 K, the carrier mobility reaches 23.4 cm2V-1s-1. At temperatures above 100 K, the carrier mobility approximately exhibits a T3/2 dependence, reflecting the dominance of ionized impurity scattering. After annealing, the carrier mobility slightly increases, thanks to a decreased defect scattering. The carrier mobility of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03), both un-annealed and annealed, exhibits a T-1/2 dependence at temperatures above 100 K, implying a scattering mechanism dominated by alloy scattering, which is consistent with the results reported in the literature

10

that although ZrNiSn system

should be a highly ordered structure, it behaved as a disordered alloy in many physical properties. The carrier mobility of both un-annealed and annealed samples of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) increases with the increasing doping content of Sb. According to the variable temperature carrier concentration for un-annealed and annealed samples of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03), the disappearance of the sharp increase of carrier concentration at around 50 K indicates the dramatic suppression of 19



the Ni interstitials concentration after doping with Sb, leading to the increase of carrier mobility. In addition, the carrier mobility of all samples increased after annealing, because the concentration of Ni interstitials decreased quite dramatically and scattering from Ni interstitial was thus diminished.

Figure 6. (a) Temperature dependent carrier concentrations in samples of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) after SHS-SPS; (b) Temperature dependent carrier mobility in samples of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) after SHS-SPS.

Figure 7 shows temperature dependent thermoelectric properties of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys prepared by the SHS-SPS method. Figure 7 (a) shows the temperature dependent electrical conductivity. The electrical conductivities of both annealed and un-annealed ZrNiSn increase with the increase of temperature in the entire temperature range investigated, displaying conduction characteristics of a semiconductor. The electrical conductivity decreases slightly after annealing, primarily because of a decreased carrier concentration. The room temperature electrical conductivity of un-annealed ZrNiSn is 1.57×104 Sm-1, comparable to that of 20


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the AM-SPS prepared sample with 1.51×104 Sm-1

14

and somewhat lower than that of

the LM-SPS prepared sample having the conductivity of 2.84×104 Sm-1

10

. The

electrical conductivity of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03), both annealed and un-annealed, decreases with the increasing temperature in the whole temperature range studied, behaving as a degenerate semiconductor. Unlike un-doped ZrNiSn, the electrical conductivity of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) slightly increases after annealing despite the decreased carrier concentrations. Obviously, the enhanced carrier mobility in this case more than compensates for the decreased carrier concentration. Figure 7(b) shows the temperature dependence of the Seebeck coefficient. The Seebeck coefficient of all samples is negative, indicating an n-type of conduction. After annealing, the Seebeck coefficient of all samples increases because the carrier concentration has decreased. The Seebeck coefficient of un-annealed ZrNiSn ranges from -240 µVK-1 to -220 µVK-1 in the whole temperature range covered. With the increasing Sb content, the Seebeck coefficient of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) alloys at 300 K decreases from -240 µVK-1 to -80 µVK-1. Figure 7(c) shows the temperature dependence of the power factor. The power factors of un-doped ZrNiSn, both annealed and un-annealed, increase with the increasing temperature in the entire measured temperature range. The power factors of Sb-doped ZrNiSn1-xSbx, both annealed and un-annealed, increase with the increasing temperature, and reach maximum power factors at around 780 K. Above 780 K, the power factors decrease with the increasing temperature. As the doping level of Sb 21



increases, the power factor increases, and the sample with the Sb content of 0.02 possesses the highest power factor. With a further Sb content increase, the power factor decreases. The maximum power factor of un-doped ZrNiSn reaches 3.6 mWm-1K-2 at 870 K, and the maximum power factor of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys attains 4.8 mWm-1K-2 at 780 K when the Sb content is 0.02 after annealing. The power factor of un-doped ZrNiSn does not changes much after annealing due to the decrease in the electrical conductivity and the concurrent increase in the absolute value of the Seebeck coefficient. The power factor of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) alloys has increased a bit after annealing because of the increase in both the electrical conductivity and the absolute value of the Seebeck coefficient. Figure 7(d) shows the temperature dependent thermal conductivity. The thermal conductivities of both annealed and un-annealed ZrNiSn decrease with the increasing temperature, and reach the minimum thermal conductivities at around 750 K. Above 750 K, the thermal conductivities increase with the increasing temperature. The thermal conductivities of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) alloys decrease with the increasing temperature in the whole temperature range. The thermal conductivity of all samples increased after annealing because of the decreased defect concentration (Ni interstitials), which significantly decreases the intensity of phonon defect scattering. With the increased content of Sb, the thermal conductivity of ZrNiSn1-xSbx (x=0.01, 0.02, 0.03) has increased from 4.8 Wm-1K-1 to 7 Wm-1K-1 at 900 K. Figure 7(e) shows the temperature dependence of the lattice thermal conductivity. 22


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The lattice thermal conductivity κL is obtained by subtracting the electronic component κe calculated via the Wiedemann-Franz law, κe = LσT, where L is the Lorenz number calculated using the SPB approximation

35

, from the total thermal

conductivity κ. At room temperature, the lattice thermal conductivity κL of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys follows the κL ~ T-1/2 dependence, indicating alloy scattering is the dominant scattering mechanism 36. As the Sb content increases, although the concentration of substitution defect of Sb on Sn site increases, it inhibits the formation of Ni interstitials which is dominated to suppress the lattice thermal conductivity. Hence, the intensity of point defect scattering from Ni interstitials is reduced leading to an increase in κL at room temperature for Sb doped ZrNiSn1-xSbx samples in comparison with un-annealed ZrNiSn sample. At high temperatures, the κL of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys follows the κL ~ T-1 dependence, indicating the dominance of Umklapp scattering of acoustic phonons 37. The κL of all samples increases after annealing due to a decreased density of Ni interstitials or defects acting as point defects. The κL of un-annealed ZrNiSn reaches 3.8 Wm-1K-1 at 870 K and the κL of un-annealed ZrNiSn0.98Sb0.02 reaches a minimum of 3.36 Wm-1K-1 at 870 K. Figure 7(f) depicts the temperature dependence of the dimensionless thermoelectric figure of merit ZT. With the increasing temperature, the ZT value increases linearly. Despite the changes in the electrical conductivity, Seebeck coefficient and thermal conductivity, the ZTs of all samples remain almost unchanged after annealing. The ZT of the un-annealed SHS-SPS prepared sample of ZrNiSn 23



reaches 0.67 at 870 K, which is comparable to or slightly higher than for the AM-SPS prepared ZrNiSn with ZT = 0.65 at 800 K ZT = 0.58 at 900 K

14

and the LM-SPS prepared sample with

10

. The highest ZT among ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03)

alloys reaches 0.7 at 870 K for the sample with the Sb content of 0.01.

Figure 7. Temperature dependent thermoelectric properties of ZrNiSn1-xSbx (x=0, 0.01, 0.02, 0.03) alloys prepared by SHS-SPS: (a) Electrical conductivity; (b) Seebeck coefficient; (c) Power Factor; (d) Thermal conductivity; (e) Lattice thermal conductivity; (f) ZT value.

4. Conclusions In this work, bulk single-phase ZrNiSn was successfully synthesized using the ultra-fast self-propagating high-temperature synthesis. Thermodynamic and kinetic parameters of the SHS process were studied, and phase transformations and 24


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microstructural evolution were revealed. A sequence of phase transformations that take place during the SHS reaction of ZrNiSn is as follows: Ni + Sn → Ni3Sn4 + Ni3Sn2

（1）(SHS)

Zr + Ni3Sn4

（2）(SHS)

→

ZrNiSn + Sn

Zr + Ni3Sn2 → ZrNi2Sn + ZrNiSn

（3）(SHS)

Zr + Sn + ZrNi2Sn

（4）(SHS)

→

ZrNiSn

ZrNiSn alloys with uniformly distributed elements, excellent compositional control and good thermoelectric properties were prepared in less than one hour. Due to the non-equilibrium nature of the SHS process, high concentration of Ni interstitials and some other defects may form in the sample, which can influence phonon and charge transport properties. Annealing of the samples minimizes the influence of Ni interstitials. The maximum thermoelectric figure of merit ZT of the SHS+SPS-processed ZrNiSn1-xSbx alloy reached 0.7 at 870 K. This study opens a new avenue for the fast, low-cost, and large-scale fabrication of half-Heusler thermoelectric materials.

Corresponding Author Email: [email protected] (X.Su), [email protected] (X.Tang)

ACKNOWLEDGMENTS The authors wish to acknowledge support from the National Basic Research Program of China (973 program) under project 2013CB632502, the Natural Science Foundation of China (Grant No. 51402222, 51172174, 51521001, 51401153), and the 25



111 Project of China (Grant No. B07040).

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