Nanoparticles at Grain Boundaries Inhibit the Phase Transformation of

Letter pubs.acs.org/NanoLett

Nanoparticles at Grain Boundaries Inhibit the Phase Transformation of Perovskite Membrane Yan Liu,†,‡ Xuefeng Zhu,*,† Mingrun Li,† Ryan P. O’Hayre,†,§ and Weishen Yang*,† †

State Key Laboratory of Catalysis, Dalian Institute of Chemical Physics, Chinese Academy of Sciences, 457 Zhongshan Road, Dalian 116023, China ‡ University of Chinese Academy of Sciences, Beijing, 100039, China S Supporting Information *

ABSTRACT: The high-energy nature of grain boundaries makes them a common source of undesirable phase transformations in polycrystalline materials. In both metals and ceramics, such grain-boundary-induced phase transformation can be a frequent cause of performance degradation. Here, we identify a new stabilization mechanism that involves inhibiting phase transformations of perovskite materials by deliberately introducing nanoparticles at the grain boundaries. The nanoparticles act as “roadblocks” that limit the diffusion of metal ions along the grain boundaries and inhibit heterogeneous nucleation and new phase formation. Ba0.5Sr0.5Co0.8Fe0.2O3−δ, a high-performance oxygen permeation and fuel cell cathode material whose commercial application has so far been impeded by phase instability, is used as an example to illustrate the inhibition action of nanoparticles toward the phase transformation. We obtain stable oxygen permeation flux at 600 °C with an unprecedented 10−1000 times increase in performance compared to previous investigations. This grain boundary stabilization method could potentially be extended to other systems that suffer from performance degradation due to a grainboundary-initiated heterogeneous nucleation phase transformations. KEYWORDS: Grain boundaries, nanoparticles, phase transformation, perovskite membrane, oxygen permeation impeded due to rapid permeation flux degradation induced by phase transformations in the LT range.10−13 Specifically, a deleterious phase transformation from cubic to the hexagonal and/or lamellar trigonal phases takes place as the temperature below 850 °C. The nature of this phase transformation process is revealed in Supporting Information Figure S1, and it indicates a time-dependent phase transformation process originating at the grain boundaries of the parent material. This time-dependent phase transformation process can be monitored in real time by determining the time-dependent oxygen permeation flux through the membrane, because the permeation takes place via oxygen ionic diffusion through grains and grain boundaries (Figure 1a). As the hexagonal and/ or lamellar trigonal phases appear at the grain boundaries, a remarkable decrease of oxygen permeation flux can be detected due to the much lower oxygen ionic conductivity of these newly nucleated grain boundary phases compared to the parent phase. Although traditional approaches to inhibit phase transformation, for example, by heavily doping stable valence state cations in perovskite lattice, can be successful, these strategies incur a significant loss of permeability.18 Therefore, researchers are facing a great challenge to inhibit phase transformation

G

rain boundaries are ubiquitous to all polycrystalline materials. Grain boundaries significantly influence most macroscopic properties of polycrystalline materials including basic mechanical,1 electronic,2 catalytic,3 ionic,4 and ferroelectric5 behaviors. They also often act as initiation sites for various materials degradation processes, including impurity segregation,6 pore formation,7 or phase transformation.8,9 Grain-boundary initiated materials degradation processes are a significant issue for many materials systems, including most perovskite-based materials used in high-temperature fuel cells and membranes. These materials are generally “perched on the edge of stability” because the same features that lead to excellent catalytic or permeation behaviors (e.g., the inclusion of multiple dopant species to induce mixed ionic and electronic conductivity, fast solid-state diffusion, and high catalytic activity) also commonly lead to rapid degradation.10−13 Ba0.5Sr0.5Co0.8Fe0.2O3−δ (BSCF), which was designed as an oxygen separation membrane14,15 and applied as fuel cell cathode and metal−air battery cathode,16,17 is a prototypical example of a high-performance perovskite that suffers from phase-transformation-induced performance degradation. In principle, the high ionic conductivity and catalytic activity toward oxygen exchange should enable BSCF to deliver excellent oxygen flux when employed as a permeation membrane even at low temperatures (550−650 °C).18−20 However, low-temperature (LT) operation is currently © XXXX American Chemical Society

Received: September 11, 2015 Revised: October 21, 2015

A

DOI: 10.1021/acs.nanolett.5b03668 Nano Lett. XXXX, XXX, XXX−XXX

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Figure 1. Oxygen permeation through and phase transformation of BSCF at 600 °C. (a) Illustration of oxygen permeation through a polycrystalline perovskite membrane. (b) Time-dependent oxygen permeation flux (JO2) for a 0.5 mm thick BSCF membrane. (c) Scanning transmission electron microscopy (STEM) image for the BSCF membrane after running for 500 h at 600 °C. “R” denotes the trigonal phase precipitated at the grain boundary. (d) EDS map analysis of the area marked in the red frame in (c). (e,f) HRTEM images for the interfacial area between the parent phase (BSCF) and the trigonal phase viewed along their [001] directions, respectively.

cubic structure with high oxygen ionic conductivity to a lowersymmetry structure with extremely poor oxygen ionic conductivity.11−13 This new low-symmetry phase can be identified by post-mortem high-resolution transmission electron microscopy (HRTEM) analysis of the spent BSCF membrane (Figure 1c−f). As shown in Figure 1c, a new phase with a width of ∼200 nm is observed at the grain boundary. This new phase is enriched with cobalt and depleted of strontium, as revealed by the energy dispersive spectrometry (EDS) analysis in Figure 1d. The BSCF parent phase maintains the cubic perovskite structure with a d100 = 0.399 nm (Figure 1e), which fits well with the unit-cell parameter obtained from X-ray diffraction (Supporting Information Figure S2). Meanwhile, the reconstructed three-dimensional reciprocal lattice of the new phase, as derived from selected area electron diffraction pattern analysis, indicates a trigonal structure (space group, R3̅m (No. 166)) with cell parameters of a = b = 0.565 nm and c = 3.576 nm (Supporting Information Figure S5). Detailed EDS analysis of this new phase suggests an elemental composition of approximately Ba0.45Sr0.07Co1.6Fe0.18O2.8. We have closely examined multiple HRTEM cross sections from spent BSCF membranes subjected to LT operation. In all cases, transformation to the trigonal second phase is observed to initiate at grain boundaries. The ability of the grain boundary to act as an activation site for the new phase has been previously verified in other BSCF studies,13 and there have been no observations of homogeneous nucleation of the trigonal phase in the grain bulk. Therefore, we deduce that the phase transformation is controlled by heterogeneous nucleation at the grain boundaries. To prevent the deleterious trigonal phase transformation, we first attempted to stabilize the perovskite parent phase by substituting small amounts (3 mol %) of various B-site cation dopants. Specifically, we examined Y3+, Zr4+, and Nb5+, a series of dopants that increase in valence charge and decrease in ionic size. While Zr4+ and Nb5+ do not

while also maintaining high permeability. Herein, we identify a new stabilization mechanism that involves inhibiting the undesirable LT phase transformation by deliberately introducing nanoparticles at grain boundaries. Using this novel approach, we demonstrate stable oxygen permeation flux (500 h) through a nanoparticle-decorated BSCF perovskite membrane with 10−1000 times higher permeation flux than previous studies at 550−650 °C. The BSCF membrane has a cubic perovskite structure with the Pm3̅m space group (No. 221) and a unit-cell parameter of 0.3981 nm (Supporting Information Figure S2) prior to operation. Oxygen permeation tests were accomplished using the setup shown in Supporting Information Figure S3. The typical degradation behavior of a standard BSCF membrane at 600 °C is shown in Figure 1b. The BSCF membrane is coated with Sm0.5Sr0.5CoO3−δ (SSC) catalyst on both surfaces before permeation testing, as prior investigations have shown that degradation associated with the surface oxygen exchange process can be eliminated by this procedure.15,19,21 Under low-temperature (LT) operation at 600 °C, we observe a dramatic 30% decrease in the oxygen permeation flux during 500 h of on-stream testing (Figure 1b). The changes in the oxygen permeation resistances over time were determined by using a previously developed permeation model (see Supporting Information Figure S4). The results indicate that the bulk resistance of the BSCF membrane increased by ∼20% over ∼110 h, whereas the oxygen exchange resistance for both surfaces remained nearly constant. Because the SSC coating eliminates degradation associated with the surface exchange process, we associate this LT degradation with a timedependent increase in the oxygen ion diffusion resistance through the BSCF membrane bulk including grain and grain boundaries (Figure 1a). A number of researchers have previously connected this decrease in oxygen permeation flux to a slow phase transformation of the membrane bulk from a B


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ionic radii of Zr4+ (0.072 nm) and Nb5+ (0.064 nm) more closely match the parent B site, and thus these dopants are fully soluble in BSCF at the 3 mol % doping level investigated. In concert with the strong flux stabilization provided by Y-doping, we observe that the trigonal phase formation is markedly suppressed in this material. More specifically, we observe that the trigonal phase is completely absent from all nanoparticlecontaining grain boundaries (Figure 2b), although it is present at the nanoparticle-free grain boundaries (Figure 2c). This exciting observation indicates that the deleterious trigonal phase transformation can be inhibited, or perhaps even fully eliminated, by deliberate nanoparticle decoration of the grain boundary regions. The phase transformation from cubic to trigonal is accompanied by significant elemental composition changes (Figure 1d and Supporting Information Figure S6) indicating a diffusional transformation mechanism in the solid phase. Longdistance diffusion of metal ions is required to form and grow the nuclei. Grain boundaries have higher energy than the bulk material and are active sites for the heterogeneous nucleation and subsequent growth of the trigonal phase. The driving force for heterogeneous nucleation at grain boundaries is related to the inherent grain boundary pressure (Pgb,in), and the corresponding energy is given by9

provide any LT stabilization effect (Supporting Information Figure S6), Y3+ does provide noticeable stabilization of the LT flux (Figure 2). The Y3+ doped membrane shows only 9% flux

Figure 2. Time-dependent oxygen permeation and phase transformation of Y-doped BSCF membranes at 600 °C. (a) Timedependent oxygen permeation flux through a 0.5 mm thick Y-doped BSCF membrane. (b) Scanning electron microscopy (SEM) image of a grain boundary with nanoparticles after the 500 h on-stream. (c) SEM image of a grain boundary without nanoparticles but a trigonal phase (denoted by “R”) after the 500 h on-stream.

degradation after 500 h on-stream testing at 600 °C (Figure 2a). Post-mortem analysis of the Y-doped BSCF membrane reveals the appearance of small nanoparticles along some of the grain boundaries in the material (Figure 2b). EDS analysis (Supporting Information Figure S7) suggests that the Y3+ dopant is only partially soluble in the BSCF bulk. In comparison, complete solid solubility is observed for the Zr4+ and Nb5+ dopants and no nanoparticles are found at the grain boundaries in these membranes (Supporting Information Figure S8). We deduce that the precipitation of second phase nanoparticles at the grain boundaries of Y-doped BSCF is driven by the significantly larger ionic size of Y3+ (0.09 nm) relative to that of Co2+/3+ (0.055−0.065 nm). In contrast, the

ΔGgb,in = −VPgb,in =

−3γV 2R g

(1)

where V is the molar volume of the material, γ is the grain boundary energy per unit area, and Rg is the grain radius. When nanoparticles decorate the grain boundaries, a Zener pinning pressure (PZ) is produced on the grains with a corresponding energy given by22

Figure 3. Nanoparticles inhibition of the phase transformation of Ce-doped BSCF perovskite. (a) Typical SEM image of grain boundaries for the asprepared Ce-doped BSCF membrane. (b) STEM image of the grain boundaries of the Ce-doped BSCF membrane after running for 500 h at 600 °C. NP denotes the nanoparticle. (c) HRTEM image and SAED pattern for the NP particle viewed along the [001] direction. (d) Time-dependent oxygen permeation of the 0.5 mm thick Ce-doped BSCF membrane at 600 °C. (e,f) HRTEM images of the “BSCF grain 1” and “BSCF grain 2” particles viewed along the axes closest to the [311] and [111] directions, respectively. C


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Figure 4. Schematic of the nanoparticles’ inhibition of phase transformation at grain boundaries. (Route a) Metal ions diffuse easily along the grain boundary and form a nucleus, and then the nucleus grows to trigonal phase at the grain boundary after long-term operation at low temperature. (Route b) Nanoparticles locate at the grain boundary, leading to a zigzag structure; the nanoparticles block the diffusion of metal ions along the grain boundary and inhibit the nucleation and growth of the trigonal phase at the grain boundary even after long-term operation at low temperature.

ΔGZ = VPZ = V

3γfV rNP

sin θ cos θ

Equations 3 and 5 suggest that the introduction of a sufficient density of nanoparticles at the grain boundary can reduce the grain boundary energy and the stepwise motion velocity to zero. The critical requirement is Rg f v/rNP > 1. For the Y-doped BSCF membrane examined in Figure 2, the nanoparticle coverage is too low, and it does not extend to all grain boundaries. Therefore, phase stabilization is only partially achieved, and the trigonal phase still develops along undecorated grain boundaries. However, improved nanoparticle coverage can be obtained by better choice of the dopant species. We have identified Ce4+ as a particularly promising dopant (Figure 3). Our choice of Ce4+ as a dopant for this nanoparticle stabilization approach is motivated by our previous observations that it is difficult to incorporate Ce4+ into Co- or Fe-based perovskite lattices because of its large ionic radius (0.087 nm) and high valence state.24,25 As shown in Figure 3, we observe that 3 mol % Ce-doped BSCF, produced via a facile one-pot synthesis method (Supporting Information Figure S9), results in the formation of a relatively widespread coverage of nanoparticles at BSCF grain boundaries (Figure 3a,b). Nanoparticle development is driven by the extremely low solubility of Ce in the BSCF lattice (less than 0.1 mol % as shown in Supporting Information Figure S10), and this point is also supported by nearly identical unit cell parameters for BSCF and Ce-doped BSCF (Supporting Information Figure S2). Structural analysis of the Ce-doped BSCF indicates that the nanoparticles are a BaCeO3 perovskite phase (Figure 3c, Supporting Information Figures S2 and S11). Nanoparticles 50−100 nm and 100−200 nm in size are observed at the grain boundaries and crystal edges, respectively, with an average interval between particles of less than 1 μm. The volume fraction of BaCeO3 nanoparticles in the membrane is only about 4%. It is too low to form a separate bulk phase. The Cedoped BSCF membrane yields an extremely stable oxygen permeation flux during 500 h of on-stream testing (Figure 3d). After 500 h of operation, no trigonal phase is observed at BSCF grain boundaries where the BaCeO3 nanoparticles are present

(2)

Where f v is the volume fraction of nanoparticles, rNP is the nanoparticle radius, and θ is the angle between the grain boundary surface and the surface at the point where the grain joins the particle. Zener pinning is known to significantly affect grain growth and recrystallization as well as recovery processes during material deformation.22,23 Here, we link it to the nucleation and growth thermodynamics and kinetics associated with new phase formation at grain boundaries. Combining eqs 1 and 2, the total free energy associated with nanoparticledecorated grain boundaries can be written as ⎛ 1 ⎞ f ΔGgb,total = 3γV ⎜⎜ − + V sin θ cos θ ⎟⎟ rNP ⎝ 2R g ⎠

(3)

9

According to classical nucleation theory, the trigonal phase in BSCF should grow via a step-growth mechanism. That is to say, the growth of nuclei is controlled by the stepwise motion along grain boundaries. This stepwise motion directly reflects the net diffusion of metal ions along the grain boundary with a resulting velocity (v) given by υ=

kmP h

(4)

where k, m, P, and h are the step distance, grain boundary migration mobility, driving pressure, and step height, respectively. Accounting for the Zener pinning effect of the grain-boundary nanoparticles on the driving pressure for the stepwise motion results in the following final expression for the stepwise motion velocity υ=

⎞ fV 3γkm ⎛⎜ 1 ⎟⎟ − sin θ cos θ h ⎜⎝ 2R g rNP ⎠

(5)

The detailed derivation and description of the above equations are provided in Supporting Information Note S1. D


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Figure 5. Asymmetric Ce-doped BSCF membrane and oxygen permeation. (a) Cross-sectional SEM image of an asymmetric Ce-doped BSCF membrane. (b) Influence of feed-side oxygen partial pressure and temperature on the permeation flux. (c) Influence of both permeation and feedside oxygen partial pressure on the permeation flux at 600 °C. P0 is 1 atm; P′O2 and P″O2 are the oxygen partial pressures of feed and permeation sides, respectively. (d) Comparison of the permeation fluxes of the asymmetric membrane (blue circle) with those of other membranes reported in literatures (refs 20 and 26−40,). ○, self-supported membranes; □, asymmetric membranes or hollow fibers. The oxygen partial pressure of the feedside was fixed at 1.0 atm, while that of the sweep-side varied in the range of 0.006−0.02 atm.

boundary mobility in the vicinity of a nanoparticle is greatly reduced compared to grain boundary mobility far from a nanoparticle, a zigzag character to the grain boundary structure develops during the high temperature membrane sintering process of the nanoparticle-decorated BSCF. In contrast, straight and smooth grain boundaries are observed in undoped, Zr-doped, and Nb-doped BSCF (Supporting Information Figure S13). As a final demonstration of the commercial applicability of the nanoparticle stabilization approach, we applied this Cedoping strategy to a thin, asymmetric membrane (Figure 5a). The asymmetric membrane consists of a ∼ 90 μm thick dense layer on top of a ∼700 μm thick porous support. The surface of the dense layer is coated with a porous 30 μm thick SSC layer to activate oxygen. The oxygen permeation flux increases rapidly with increasing oxygen partial pressure on the feed side (Figure 5b) or decreasing oxygen partial pressure on the permeation side (Figure 5c) and reaches approximately 3 mL cm−2 min−1 when the feed side oxygen partial pressure is 0.21 atm and reaches up to 5.2 mL cm−2 min−1 when the oxygen partial pressure is 1.0 atm at 600 °C. As temperature decreased to 550 °C, the oxygen permeation fluxes are 1.4 and 2.5 mL cm−2 min−1 when the oxygen partial pressures of feed side are 0.21 and 1.0 atm, respectively. The asymmetric membrane fluxes at temperature range of 550−650 °C are 10−1000 times higher than those reported for other membranes (Figure 5d) and exceed the commonly cited threshold (>1 mL cm−2 min−1)18 required for commercial viability as the membrane operated under oxygen partial pressure gradients of 1.0 atm/ 0.006−0.02 atm.

(Figure 3b,e,f). The oxygen permeation resistances across the Ce-doped BSCF membrane obtained through model investigation shows that the bulk resistance keeps constant (Supporting Information Figure S12), which also indicates there is no phase transformation taking place when BaCeO3 nanoparticles appear at the grain boundaries. Because the BSCF and Ce-doped BSCF membranes have been coated by the same catalyst, the similar exchange resistances of the two membranes, as shown in Supporting Information Figures S4 and S12, indicate the high reliability of the model investigation to identify the bulk changes. The nanoparticle-based stabilization effect is summarized in Figure 4. The net cationic diffusion flux for the original material at the grain boundaries aids heterogeneous nucleation and growth of the trigonal phase (route (a)). However, nanoparticles located at the grain boundaries inhibit the net diffusion flux by limiting stepwise motion along the grain boundaries (route (b)). For the Cedoped BSCF membrane presented in Figure 3, the average BSCF grain diameter is approximately 15 μm while the average nanoparticle diameter is approximately 100 nm; the nanoparticle volume fraction is about 4%, yielding a Rg f v/rNP value of approximately 6. This value is much larger than the critical value of 1 and indicates, theoretically, that the driving force for heterogeneous nucleation and growth of the trigonal phase via stepwise motion has been reduced to zero. In other words, the nanoparticles are successfully acting as “roadblocks” and have reduced the net diffusion flux of cations along the nanoparticledecorated grain boundaries to zero. Inhibition of the stepwise motion mechanism is also supported by the zigzag nature of the grain boundaries seen in the Ce-doped material. Because grain E


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(14) Shao, Z.; Yang, W.; Cong, Y.; Dong, H.; Tong, J.; Xiong, G. J. Membr. Sci. 2000, 172, 177−188. (15) Liu, Y.; Zhu, X.; Yang, W. AIChE J. 2015, 61, 3879−3888. (16) Shao, Z.; Haile, S. M. Nature 2004, 431, 170−173. (17) Suntivich, J.; May, K. J.; Gasteiger, H. A.; Goodenough, J. B.; Shao-Horn, Y. Science 2011, 334, 1383−1385. (18) Sunarso, J.; Baumann, S.; Serra, J. M.; Meulenberg, W. A.; Liu, S.; Lin, Y. S.; Diniz da Costa, J. C. J. Membr. Sci. 2008, 320, 13−41. (19) Liu, Y.; Zhu, X.; Li, M.; Liu, H.; Cong, Y.; Yang, W. Angew. Chem., Int. Ed. 2013, 52, 3232−3236. (20) Tong, J.; Yang, W.; Zhu, B.; Cai, R. J. Membr. Sci. 2002, 203, 175−189. (21) Liu, Y.; Zhu, X.; Li, M.; Li, W.; Yang, W. J. Membr. Sci. 2015, 492, 173−180. (22) Nes, E.; Ryum, N.; Hunderi, O. Acta Metall. 1985, 33, 11−22. (23) Humphreys, F. J. Acta Mater. 1997, 45, 5031−5039. (24) Zhu, X.; Wang, H.; Yang, W. Solid State Ionics 2006, 177, 2917− 2921. (25) Li, Q.; Zhu, X.; Yang, W. Mater. Res. Bull. 2010, 45, 1112−1117. (26) Gaudillere, C.; Garcia-Fayos, J.; Serra, J. M. J. Mater. Chem. A 2014, 2, 3828−3833. (27) Girdauskaite, E.; Ullmann, H.; Vashook, V.; Guth, U.; Caraman, G.; Bucher, E.; Sitte, W. Solid State Ionics 2008, 179, 385−392. (28) Haworth, P.; Smart, S.; Glasscock, J.; Diniz da Costa, J. C. Sep. Purif. Technol. 2012, 94, 16−22. (29) Jiang, Q.; Faraji, S.; Nordheden, K. J.; Stagg-Williams, S. M. J. Membr. Sci. 2011, 368, 69−77. (30) Kida, T.; Takauchi, D.; Watanabe, K.; Yuasa, M.; Shimanoe, K.; Teraoka, Y.; Yamazoe, N. J. Electrochem. Soc. 2009, 156, E187−E191. (31) Rachadel, P. L.; Motuzas, J.; Ji, G.; Hotza, D.; Diniz da Costa, J. C. J. Membr. Sci. 2014, 454, 382−389. (32) Serra, J. M.; Garcia-Fayos, J.; Baumann, S.; Schulze-Küppers, F.; Meulenberg, W. A. J. Membr. Sci. 2013, 447, 297−305. (33) Kharton, V. V.; Yaremchenko, A. A.; Kovalevsky, A. V.; Viskup, A. P.; Naumovich, E. N.; Kerko, P. F. J. Membr. Sci. 1999, 163, 307− 317. (34) Wang, H.; Tablet, C.; Caro, J. J. Membr. Sci. 2008, 322, 214− 217. (35) Watenabe, K.; Yuasa, M.; Kida, T.; Teraoka, Y.; Yamazoe, N.; Shimanoe, K. Adv. Mater. 2010, 22, 2367−2370. (36) Wu, Z.; Hidayati Othman, N.; Zhang, G.; Liu, Z.; Jin, W.; Li, K. J. Membr. Sci. 2013, 442, 1−7. (37) Qi, X.; Lin, Y. S.; Swartz, S. L. Ind. Eng. Chem. Res. 2000, 39, 646−653. (38) Zhang, G.; Liu, Z.; Zhu, N.; Jiang, W.; Dong, X.; Jin, W. J. Membr. Sci. 2012, 405−406, 300−309. (39) Wang, Z.; Yang, N.; Meng, B.; Tan, X. Ind. Eng. Chem. Res. 2009, 48, 510−516. (40) Shao, Z.; Xiong, G.; Tong, J.; Dong, H.; Yang, W. Sep. Purif. Technol. 2001, 25, 419−429.

In summary, we have demonstrated that nanoparticles at the grain boundaries can effectively inhibit the heterogeneous nucleation and growth of the trigonal phase in BSCF. This approach enables high stability and high oxygen permeability to be simultaneously achieved in the LT regime. While we demonstrate this grain boundary stabilization method for BSCF in particular, we suggest that it could potentially be extended to other systems that suffer from performance degradation due to a grain-boundary-initiated heterogeneous nucleation phase transformations.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.5b03668. Experimental details and additional figures. (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. Present Address §

(R.O.) Metallurgical and Materials Engineering Department, Colorado School of Mines, 1500 Illinois Street, Golden, Colorado 80401-1887, United States. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS All authors thank the financial support from Natural Science Foundation of China (21271169 and 21476225) and the Key Research Program of the Chinese Academy of Sciences (Grant KGZD-EW-T05). R.O. thanks the Chinese Academy of Sciences Visiting Professorships (Grant 2012T1G0015).

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Nanoparticles at Grain Boundaries Inhibit the Phase Transformation of

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