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Article

Near-Zero Thermal Expansion and High Strength in Multi-Phase La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag Compounds Produced Through Spark Plasma Sintering

by
Yuyu Wang
1,2,†,
Kai Xu
1,2,†,
Hanyang Qian
3,
Rui Cai
3,
Xiang Lu
3 and
Jian Liu
1,2,*
1
State Key Laboratory of Materials for Advanced Nuclear Energy, Shanghai University, Shanghai 200444, China
2
School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China
3
Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, China
*
Author to whom correspondence should be addressed.
†
These authors contributed equally to this work.
Metals 2025, 15(10), 1131; https://doi.org/10.3390/met15101131 (registering DOI)
Submission received: 21 July 2025 / Revised: 11 September 2025 / Accepted: 19 September 2025 / Published: 11 October 2025
(This article belongs to the Section Metallic Functional Materials)
Browse Figures
Versions Notes

Abstract

The significant negative thermal expansion (NTE) that occurs in La(Fe,Si)13-based alloys during magnetic transition make them promising to combine with positive thermal expansion (PTE) materials to obtain near-zero thermal expansion (NZTE) materials. However, La(Fe,Si)13-based alloys with NTE generally show intrinsic poor mechanical properties. Here, thermal expansion properties are optimized by adding Ag in La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1 to form a multi-phase structure exhibiting enhanced compressive strength. Through spark plasma sintering (SPS) and annealing, the samples consisted of α-Fe(Co,Si), NaZn13-type, and LaAg2 phases. When the annealing temperature reaches 1323 K, LaAg2 disappears and is replaced by La2O3. The LaAg2 phase and α-Fe(Co,Si) phase contribute as PTE materials to compensate for the NTE of the NaZn13-type phase. Near-zero thermal expansion was achieved in the temperature range of 240–294 K, with a coefficient of thermal expansion (CTE) of 3.5 ppm/K at a 9.581 at.% Ag content. Benefiting from the uniform phase distribution and coordinated deformation, the samples obtained high mechanical strengths, with fracture stresses of 1481.1 MPa for the 15 wt.% Ag sample. This work provides a promising route for high-strength and near-zero thermal expansion Ag/La(Fe,Si)13 composites.

1. Introduction

Most substances expand when heated and contract when cooled, which is described as positive thermal expansion (PTE). Thermal fluctuations cause almost negligible dimensional changes in most solid materials, but in specific fields, such as aerospace and precision sensors, they can lead to functional failure of critical components. In recent years, numerous materials demonstrating contraction upon heating and expansion upon cooling have been identified, classified as negative thermal expansion (NTE) materials [1]. NTE materials facilitate precise adjustment of the coefficient of thermal expansion (CTE) in composite systems to obtain zero thermal expansion materials suitable for high-precision devices [2,3,4,5,6,7,8,9,10]. Diverse NTE materials exhibiting tunable CTE have been reported, including the ZrW2O8, MnCoGe-based alloys [11,12], PbTiO3-based compounds [13,14], SrCu3Fe4O12 [15], Bi0.95La0.05NiO3 [16], and magneto-volume effect materials, such as Hf(Ta,Fe)2 [17,18,19,20] and La(Fe,Si)13 [21,22,23,24,25,26].
La(Fe,Si)13-based alloy, characterized by the NaZn13-type cubic structure, represents a prototypical category of NTE materials driven by magnetic phase transitions [27,28]. A structural transition from the ferromagnetic to paramagnetic state occurs near the Curie point with increasing temperature [25,26]. This magneto-volume effect induces substantial lattice contraction, manifesting as pronounced NTE behavior [22]. The La(Fe,Si)13 system exhibits isotropic NTE due to the highly symmetrical structure and the magnetoelastic transition of isotropic lattice contraction [18,23]. However, a common limitation of magnetostrictive NTE materials is their poor mechanical properties. This seriously affects their utility as NTE components subjected to cyclic thermal loading.
A composite of La(Fe,Si)13 alloy with the second phase is a double-win method to improve the mechanical strength and adjust the CTE. Peng et al. [29] achieved a compressive strength of 240.8 MPa by 20 wt.% Cu composite with La(Fe,Si)13. In La(Fe,Si)13/Cu composite, benefiting from the weakening of the magneto-volume effect by Cu, a near-zero thermal expansion is shown when the Cu content is 45 wt.% [30]. In addition, Al [31], In [32], Bi [33], and Pr2Co7 [34] have also been used to fabricate composites to improve the mechanical properties of La-Fe-Si alloys. However, this method failed to establish a strong interfacial bond between the matrix and the second phase, thereby substantially enhancing the mechanical strength [32,33,35,36,37,38]. In addition, the introduction of an in situ precipitated α-Fe phase to compensate for the negative thermal expansion of the NaZn13-type phase can achieve NZTE and enhance mechanical properties [39,40]. Shao et al. [41] showed that the presence of parallel orientations of the La(Fe,Si)13 phase and the α-Fe(Co,Si) phase at the phase interface is the key to the enhancement of thermal conductivity and strength. In addition, different preparation methods have a strong influence on the strength of La(Fe,Si)13-based alloys. Zhong et al. [42] prepared La-Fe-Si alloys through SPS and hot-pressing sintering methods. The fracture strength of the SPS sample was 637.9 MPa, which was significantly higher than that of the hot-pressing sintering sample (275.4 MPa). Therefore, SPS is frequently used as a powerful approach for the preparation of composite materials.
In this work, La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag alloys were prepared using the SPS method. The addition of Ag significantly enhanced the mechanical properties of the alloys. The annealed samples consisted of the NaZn13-type phase, α-Fe(Co,Si), and LaAg2 phases. The NaZn13-type phase exhibited negative thermal expansion, while the other phases exhibited positive thermal expansion, and compensation of negative and positive thermal expansion resulted in tunable negative to near-zero thermal expansion by adjusting the Ag content.

2. Materials and Methods

Gas-atomized powder containing 6.17 La, 4.11 Ce, 71.97 Fe, 7.12 Co, and 10.63 Si (in at.%) was used in this work. It was uniformly mixed with Ag spherical powder (average particle size of 10 μm) in 2.852 at.% and 9.851 at.%, respectively. Figure 1a shows the SEM image of the mixed powders, which have good sphericity. The average particle size of the mix was 7.55 ± 0.13 μm (Figure 1b). Figure 1c shows the EDS mapping of the mixed powders. As displayed in Figure 1c, the elemental mapping profile obtained from EDS reveals uniform distributions of the elements La, Ce, Fe, Co, Si, and Ag. The La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/2.852 at.% Ag, La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/9.851 at.% Ag composite alloys obtained through SPS at 40 MPa, 1123 K for 10 min were defined as S1 and S2, respectively. The S1 and S2 samples annealed at 1273 K for 6 h were defined as HT-S1 and HT-S2, respectively.
The microstructure and chemical composition were analyzed using an electron probe microanalyzer (EPMA, SHIMADZU 8050G, Shimadzu, Kyoto, Japan) in backscattered electron (BSE) mode and an energy dispersive spectrometer (EDS, OXFOED ULTIMMAX655, Oxford Instruments, Abingdon, UK). The powder morphology was observed through scanning electron microscopy (SEM, ZEISS Sigma 300, ZEISS, Oberkochen, Germany). The phase constitution and crystal structure of the samples were determined through X-ray diffraction (XRD, Malvern Panalytical Empyrean, Cu Kα radiation, Malvern Panalytical, Malvern, UK) and refined with the Rietveld method using FullProf software (Version January-2021). The room temperature mechanical properties were obtained through mechanical compression performed on a 3 × 3 × 4 mm3 sample using an MTS CMT5202 universal testing machine (MTS Systems, Eden Prairie MN, USA). A DIC (Correlation Solutions Inc., Irmo, SC, USA) system was used to investigate the thermal expansion behavior of HT-S1 and HT-S2 samples. The temperature dependence of the magnetization was measured using a superconducting quantum interference device (SQUID, Quantum Design, San Diego, CA, USA). The DIC camera had a resolution of 4906 × 2160 pixels, with speckle sizes ranging from 0.1 to 0.8 mm (average: 0.2 mm). Images were captured at 1 Hz and processed using commercial software VIC-2D (Digital lmage Correlation Version 6.0.6, build 665 Copyright 1998–2017 Correlated Solutions, Inc.). The linear thermal expansion (ΔL/L0) was calculated using virtual extensometers.

3. Results and Discussion

3.1. Microstructural Evolution

Figure 2 shows the SEM images of the sintered samples; no pores or cracks were observed, and high densification was obtained. The Ag phase predominantly filled the interstices between spherical powder particles. The spherical powder reduces the irregular deformation of the particles caused by axial pressure during sintering. Compared with irregular powders, spherical powders have fewer sintering necks, which is beneficial for accelerating elemental diffusion and improving densification [42]. Figure 2a,c show SEM images of S1 and S2 samples, respectively. The microstructures of both samples are similar, and the complete La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1 spherical powder can still be observed after sintering. Figure 2b,d show magnified images of the areas marked in yellow in Figure 2a,c, respectively. The sintered sample consists of spherical particles and a white contrast phase filled between the particles. Two phases with different contrasts can also be observed in the spherical powder. One is a dark contrasting dendritic phase, and the other is a white contrasting phase between the dendritic phases. Different from common two-phase structures, the addition of Ag causes changes in the phase composition of the alloy. Figure 2e shows the XRD pattern of the S1 sample. The sintered sample is mainly composed of α-Fe(Co,Si), CeFe2Si2, LaCo2Si2, and La14Ag51 phases, and no (La,Ce)(Fe,Co)Si phase was identified. Figure 2f shows the La-Ag binary phase diagram [43]. La14Ag51 is easily formed when the Ag content is around 80%. During sintering, La diffuses from the (LaCe)(Fe,Co)Si phase into Ag to form the La-Ag phase, which is driven by thermal and stress fields. The enthalpy of formation of La14Ag51 is lower than that of (LaCe)(Fe,Co)Si [44]. Owing to the low La content, the original (LaCe)(Fe,Co)Si phase decomposes into the CeFe2Si2 and LaCo2Si2 phases.
To reveal the elemental diffusion during sintering, an EMPA mapping image of the S1 sample is shown in Figure 3. In the white contrast region between the spherical particles, there is a segregation of Ag, La, Ce, Co, and Si. This implies that during the sintering process, La, Ce, Co, and Si elements diffuse from the spherical particles to this region and form a multi-phase structure. This agrees with the phase composition identified in Figure 2f. The small amount of oxygen observed is attributed to the oxidation of rare earth elements. In addition, Ag elements diffuse into the interior of the spherical particles, and the extent of diffusion depends on the particle size. The Fe remained in the inner part of the spherical powder, and no diffusion was observed. This result is consistent with the presence of a significant amount of α-Fe(Co,Si) in the XRD results.

3.2. The Microstructure After Annealing

Figure 4a,c show the SEM images of HT-S1 and HT-S2 samples, respectively. Compared with the HT-S1 sample, the HT-S2 sample has more white contrast phases. In addition, for the HT-S1 sample, there are some punctate black phases present in the gray matrix. Figure 4b,d show magnified images of the regions marked with yellow frames in Figure 4a,c, respectively. After annealing, both of them have basically the same phase composition. Light gray, dark gray, and white contrast regions can be observed in Figure 4b. The light gray contrast phase contains a small amount of dot-like dark gray phase. Additionally, a dendritic two-phase microstructure composed of dark gray and white contrast phases is present. Figure 4e shows the EPMA mapping results for the HT-S1 sample. La and Ag segregation is observed in the white contrast region, with a small amount of oxides present. The segregation of Ce is distinct from La, caused by the enthalpy of formation with Ag, which makes it difficult to form Ag-Ce compounds in this region [45,46]. The dark gray contrast region exhibits Fe segregation, composed of the α-Fe(Co,Si) phase. The light gray contrast phase is the NaZn13-type phase.
To support phase evolution during the annealing process, the XRD patterns of the S1 sample were measured after annealing at different temperatures for 6 h (Figure 5). For the samples annealed at 1173 K for 6 h, the α-Fe(Co,Si), NaZn13-type phase and the LaAg2 phase can be identified in the XRD pattern. This corresponds to the three different contrasting phases observed in the SEM image. Compared to the unannealed sample (Figure 2e), the complexity of the phase composition is reduced. The annealing process facilitates the reaction involving α-Fe(Co,Si), CeFe2Si2, and LaCo2Si2 phases to form the NaZn13-type phase. Meanwhile, a part of the La element continues to combine with Ag to form the LaAg2 phase. This part of La combines with Ag, resulting in an incomplete peritectic reaction. Thus, a significant amount of dendritic α-Fe(Co,Si) can be observed in the annealed samples. Notably, the annealing resulted in the formation of a PTE/NTE composite, in which the α-Fe(Co,Si), LaAg2 is responsible for the PTE and the NaZn13-type phase is responsible for the NTE.
Figure 5 shows the XRD patterns of the S1 samples annealed for 6 h at different temperatures after Rietveld refinement. A three-phase refinement was performed using the FullProf software package to obtain accurate phase constitution. Table 1 summarizes the phase compositions and mass occupancies in samples annealed at different temperatures. For 1223 K and 1273 K, the annealed samples consist of the NaZn13-type phase, the α-Fe(Co,Si) phase, and the LaAg2 phase. For 1323 K, the sample consists of the NaZn13-type phase, the α-Fe(Co,Si) phase, and the La2O3 phase. There is less change in the amount of the NaZn13-type phase in samples when annealing at different temperatures. The content of the LaAg2 phase decreases as the annealing temperature increases, and the LaAg2 phase disappears and La2O3 appears when the temperature rises to 1323 K. At 1273 K, the composition of the HT-S2 sample and the HT-S1 phase is basically the same. The HT-S2 sample contains more Ag and generates more LaAg2. Due to the formation of more compound LaAg2, the 1:13 phase content decreases and the α-Fe content increases.

3.3. Thermal Expansion and Mechanical Properties

The paramagnetic/ferromagnetic transition of La(Fe,Si)13 compounds near Tc is accompanied by large magneto-volume effects [47,48]. The magnetic ordering is broken when the temperature is increased to Tc. When itinerant electronic systems need to absorb energy greater than that of thermal fluctuations, the lattice will contract to compensate for the extra demand, resulting in a large NTE effect [21]. Figure 6a,b show the linear thermal expansion ΔL/L (reference temperature: 180 K) versus temperature for the HT-S1 and HT-S2 samples in the temperature range of 180–320 K, respectively. For the HT-S1 sample, at temperatures below 228 K, ΔL/L increases as the temperature increases, showing PTE. Because there is no magnetic transition in this temperature range, the lattice expansion derives from thermally driven anharmonic atomic vibrations [4,49]. In the range of 228–289 K, ΔL/L decreases, and the coefficient of thermal expansion (CTE = ΔL/(L × ΔT)) is −11.68 × 10−6 K−1. In this temperature range, a magnetic transition occurs, leading to an NTE, and the contraction of the NaZn13-type phase is sufficient to compensate for the positive thermal expansion of LaAg2 and α-Fe(Co,Si). Above 289 K, anharmonic atomic vibrations dominate, and PTE occurs again. This coincides with the strain distribution of the HT-S1 sample during heating (Figure 6c). When the temperature reaches 228.1 K, NTE starts to appear in some regions (blue color), and the proportion of NTE regions increases with increasing temperature. At 291.2 K, a more continuous NTE region and a small amount of PTE region can be observed, and the distribution is uniform. This indicates that the PTE/NTE composites have better coordinated deformation, which is beneficial for reducing the deformation stress and improving the bond strength.
For the HT-S2 sample, PTE is shown in the temperature range of 180–320 K. It has a lower CTE = 3.25 × 10−6 K−1 in the temperature range of 240–294 K. The lattice contraction of the NaZn13-type phase induced by the magnetic transition has offset most of the PTE of LaAg2 and α-Fe(Co,Si). Compared to the annealed S1 sample, the higher content of Ag-rich phases results in the PTE not being fully compensated. Figure 6d confirms the above analysis, with a significant increase in the fraction of PTE regions (red and yellow colors) and a decrease in the fraction of NTE regions (blue and pink colors) compared to the HT-S1 sample. This result also corresponds to the observation in Figure 4b,d that more LaAg2 and α-Fe(Co,Si) phases are present in the HT-S2 sample.
Figure 7 shows the M–T curves of the HT-S1 sample. The conclusion temperature of the magnetic transformation coincides with that of the negative thermal expansion (NTE). A discrepancy is observed, however, between the onset temperatures of the NTE and the magnetic transition, which is attributed to differences in experimental conditions—specifically, the use of infrared thermometry versus PPMS measurements, as well as differing heating and cooling rates.
Except for a low CTE, good mechanical strength is a premise for maintaining long-term service stability. Figure 8a shows the compressive stress–strain curves of HT-S1 and HT-S2 samples, both of which show complete brittle fracture modes. The compressive fracture stress of the HT-S1 sample is 1269.4 MPa with a strain of 1.86%, while the fracture stress of the HT-S2 is 1481.1 MPa with a strain of 2.57%. Notably, in the stoichiometric La(Fe,Si)13 alloy, the fracture strength is much lower than in this work, and the fracture strain is usually less than 1% [30,31]. The high brittleness and low plasticity are attributed to the large differences in atomic diameters in the lattice and the lack of removable dislocations [50]. The introduction and uniform distribution of the LaAg2 phase significantly increased the mechanical strength of the alloy. The higher fracture strength of the HT-S2 sample compared to the HT-S1 sample is ascribed to the higher content of the LaAg2 phase (Figure 4a,c). Figure 8b summarizes the CTE and compressive fracture stresses of common near-ZTE materials. For ZnSnMn/Zn composites, ZnSnMn corresponds to NTE, and Zn is used as a compensating material. The increase in Zn content leads to a change from NTE to PTE, and near-ZTE can be obtained with a 70 vol.% Zn in the temperature range of 253–304 K [51]. However, the mechanical strength limits application in high-stress environments. For Zr55Cu30Ni5Al10/beta-LiAlSiO4 composite, increasing the LiAlSiO4 (PTE) decreases the mechanical strength. At LiAlSiO4 contents in the range of 20–60 vol.%, low CTE was observed, and the fracture stress was below 700 MPa [52]. For the Hf0.87Ta0.13Fe2Cux (0.1 ≤ x ≤ 1.5) alloy, Cu acts not only as PTE compensation, as its embedding in the C14 laves phase enhances the mechanical properties (1235 MPa) [20]. In addition, in La(Fe,Si)13/(Cu/Fe) composite, the excess Fe favors the compressive strength [39], and Cu acts as a ductile material providing additional plasticity and thermal conductivity [30]. In comparing the different systems of NZTE materials, the La(Fe,Si)13-based system has a narrower applicable temperature, limited by the magnetic transition temperature range. Compared to the above existing NZTE systems, HT-S1 and HT-S2 have better mechanical strength and a wider temperature range compared to the first-order magnetic transition La(Fe,Si)13 system.
Figure 9 presents the fracture surface of the HT-S1. The cleavage planes, river patterns, and a small number of dimples can be observed. This morphology indicates that the predominant fracture mode is brittle, consistent with the stress–strain curve. The dimples are associated with the presence of Ag particles and the α-Fe phase. However, due to their low volume fraction, these ductile features do not alter the dominant brittle fracture mechanism. In composite materials, second-phase toughening mechanisms generally contribute to an improvement in fracture strength.

4. Conclusions

This work investigates the microstructure, thermal expansion, and mechanical properties of La-Ce-Fe-Co-Si/Ag multi-phase alloys fabricated through SPS. The results show that Ag readily forms the La14Ag51 phase with La after sintering at 1073 K for 10 min. The sintered sample consists of α-Fe(Co,Si), CeFe2Si2, LaCo2Si2, and La14Ag51 phases. After annealing at 1273 K for 6 h, the sample consisted of α-Fe(Co,Si), NaZn13-type, and LaAg2 phases. The LaAg2 and α-Fe(Co,Si) phases were applied as PTE materials to compensate for the NTE of the NaZn13-type phase. For the HT-S1 sample, the linear CTE is −11.68 ppm/K, and the fracture stress is 1249.6 MPa. HT-S2 shows a near-zero thermal expansion with a CTE of 3.25 ppm/K and a fracture stress of 1489.8 MPa. The high mechanical strengths are attributed to the uniformly dispersed α-Fe(Co,Si) and LaAg2 phases. This work reveals that the introduction of extra Ag and the formation of NTE/PTE multi-phase composites is an effective way to improve the mechanical properties of La(Fe,Si)13 alloys and will facilitate their application as NTE materials in high-stress environments.

Author Contributions

Conceptualization, X.L. and J.L.; methodology, J.L.; software, Y.W.; validation, K.X., X.L., and J.L.; formal analysis, Y.W., K.X., H.Q., R.C., X.L., and J.L.; investigation, Y.W., H.Q., and R.C.; resources, J.L.; data curation, Y.W., H.Q., and R.C.; writing—original draft, Y.W., K.X., H.Q., and R.C.; writing—review and editing, K.X., X.L., and J.L.; visualization, Y.W.; supervision, J.L.; project administration, J.L.; funding acquisition, J.L. All authors have read and agreed to the published version of the manuscript.

Funding

This work is supported by the Shandong Provincial Key Research and Development Program (Major Science and Technology Innovation Project) (2023CXGC010301), the National Natural Science Foundation of China (52371192, 52201234), and the Natural Science Foundation of Ningbo (2022J309).

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflict of interest.

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Metals 15 01131 g001
Figure 1. (a) SEM images of mixed powders. (b) Particle size distribution. (c) The backscattered electron and EDS images of Ag and La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1 spherical particles.
Figure 1. (a) SEM images of mixed powders. (b) Particle size distribution. (c) The backscattered electron and EDS images of Ag and La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1 spherical particles.
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Figure 2. Backscattered electron (BSE) images of (a,b) S1 sample and (c,d) S2 sample, (e) XRD pattern of S1 sample, (f) Ag-La binary phase diagram, adapted from Ref. [43].
Figure 2. Backscattered electron (BSE) images of (a,b) S1 sample and (c,d) S2 sample, (e) XRD pattern of S1 sample, (f) Ag-La binary phase diagram, adapted from Ref. [43].
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Figure 3. EPMA mapping images of the S1 sample.
Figure 3. EPMA mapping images of the S1 sample.
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Figure 4. (a,b) BSE images of HT-S1 samples, (c,d) BSE images of HT-S2 samples, (e) EPMA mapping results of HT-S1 sample.
Figure 4. (a,b) BSE images of HT-S1 samples, (c,d) BSE images of HT-S2 samples, (e) EPMA mapping results of HT-S1 sample.
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Figure 5. The Rietveld refined XRD pattern of the S1 sample annealing at different temperature for 6 h. Red symbols: experimental data. Black line: calculated pattern. Blue line: the difference between experimental and calculated patterns.
Figure 5. The Rietveld refined XRD pattern of the S1 sample annealing at different temperature for 6 h. Red symbols: experimental data. Black line: calculated pattern. Blue line: the difference between experimental and calculated patterns.
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Figure 6. (a) Temperature dependence of linear thermal expansions ΔL/L (reference temperature: 180 K) of HT-S1 sample, (b) temperature dependence of linear thermal expansions ΔL/L (reference temperature: 180 K) of HT-S2 sample, (c) strain distribution in HT-S1 sample, (d) strain distribution in HT-S2 sample.
Figure 6. (a) Temperature dependence of linear thermal expansions ΔL/L (reference temperature: 180 K) of HT-S1 sample, (b) temperature dependence of linear thermal expansions ΔL/L (reference temperature: 180 K) of HT-S2 sample, (c) strain distribution in HT-S1 sample, (d) strain distribution in HT-S2 sample.
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Figure 7. The M-T curves of the HT-S1 sample annealing at 1273 K.
Figure 7. The M-T curves of the HT-S1 sample annealing at 1273 K.
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Figure 8. (a) The compressive stress–strain curves of HT-S1 and HT-S2 samples, (b) summary of CTE-compressive stress correlation in recent studies for samples annealed at 1000 °C.
Figure 8. (a) The compressive stress–strain curves of HT-S1 and HT-S2 samples, (b) summary of CTE-compressive stress correlation in recent studies for samples annealed at 1000 °C.
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Figure 9. The SEM images of the fracture surface for the HT-S1 sample (a), and an enlarged view (b).
Figure 9. The SEM images of the fracture surface for the HT-S1 sample (a), and an enlarged view (b).
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Table 1. Phase occupancy and fitting coefficients of HT-S1 samples after annealing at different temperatures.
Table 1. Phase occupancy and fitting coefficients of HT-S1 samples after annealing at different temperatures.
SampleT (K)Weight Percentage (wt.%)Fit Coefficient
NaZn13α-Fe(Co,Si)LaAg2La2O3RpRwp
HT-S1122361.0533.585.37ND3.434.38
127362.3733.973.65ND3.254.18
132360.8735.17ND3.963.464.36
HT-S2127353.3338.997.67ND4.295.38
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MDPI and ACS Style

Wang, Y.; Xu, K.; Qian, H.; Cai, R.; Lu, X.; Liu, J. Near-Zero Thermal Expansion and High Strength in Multi-Phase La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag Compounds Produced Through Spark Plasma Sintering. Metals 2025, 15, 1131. https://doi.org/10.3390/met15101131

AMA Style

Wang Y, Xu K, Qian H, Cai R, Lu X, Liu J. Near-Zero Thermal Expansion and High Strength in Multi-Phase La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag Compounds Produced Through Spark Plasma Sintering. Metals. 2025; 15(10):1131. https://doi.org/10.3390/met15101131

Chicago/Turabian Style

Wang, Yuyu, Kai Xu, Hanyang Qian, Rui Cai, Xiang Lu, and Jian Liu. 2025. "Near-Zero Thermal Expansion and High Strength in Multi-Phase La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag Compounds Produced Through Spark Plasma Sintering" Metals 15, no. 10: 1131. https://doi.org/10.3390/met15101131

APA Style

Wang, Y., Xu, K., Qian, H., Cai, R., Lu, X., & Liu, J. (2025). Near-Zero Thermal Expansion and High Strength in Multi-Phase La0.6Ce0.4(Fe0.91Co0.09)11.9Si1.1/Ag Compounds Produced Through Spark Plasma Sintering. Metals, 15(10), 1131. https://doi.org/10.3390/met15101131

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