Dealloying-induced phase transformation in Fe–Rh alloys

Nanoporous metals produced by dealloying have aroused enormous interest due to exotic mechanical and physico-chemical properties that are usually inaccessible in their bulk form. Interestingly, when binary solid-solution alloys, such as Ag–Au alloys, are dealloyed, the resulting nanoporous metals usually inherit the crystal structure of their parent alloys. In this Letter, we examined the evolution of the crystal structure during the dealloying of Fe–Rh alloys that show single-phase solubility over a large range of compositions. In situ x-ray diffraction shows that the crystallographic structure of the Fe 85 Rh 15 alloy transforms from the original bcc to fcc structure during the dealloying. Transmission electron microscopy conﬁrms the fcc structure of the nanoporous sample, which exhibits a typical bi-continuous porous structure with ligament sizes of only 2–3nm and a high Fe concentration. The bcc–fcc transformation is driven by the chemical disordering of Fe and Rh atoms, induced by the highly dynamic dissolution and diffusion process at the alloy/electrolyte interface. Our study highlights the massive diffusion and the consequent disordered arrangement of elemental components during the evolution of the nanoporous structure.

Dealloying is a corrosion process that selectively dissolves one or more elemental components with relatively low electronegativity in an alloy.It has recently emerged as a powerful approach for fabricating nanoporous (NP) metals, which exhibit an open porous morphology with inter-penetrating ligaments and pores at the nanometer scale. 1,2ue to their extremely large surface area, they can feature mechanical and physico-chemical properties that are inaccessible in bulk metals such as ultrahigh specific strength, 3,4 "piezo-electric" effect, 5,6 responsive strength, 7,8 and high catalytic activity. 9,10With their overall monolithic form, NP metals also offer an extremely versatile platform to unveil hidden physical properties resulting from the strong surface effects. 11,124][15] As exemplified in the prototypical Ag-Au alloy, Ag atoms are selectively dissolved from atomic terraces, and the Au atoms remain during the dealloying, which produces the supersaturated solution of Au ad-atoms and vacancies at the alloy/electrolyte interface.Being thermodynamically unstable, this supersaturated solution decomposes into a two-phase structure of Au clusters and pores through uphill diffusion, leading eventually to the porous structure. 13n the other hand, from the experimental point of view, substantial efforts have been devoted to characterize the morphology of the resultant NP structure, especially with regard to the ligament sizes and contents of residual Ag.The observed ligament sizes span from a few nanometers to sub-micron scale 16,17 and the residual Ag concentration from 0 to 70 at.%. 18,19 This large-range scale arises from the secondary coarsening/dealloying process that occurs behind the corrosion front, which can complicate the direct comparison of experimental results with the theoretical predictions.
The understanding of the nanoporosity formation requires, among others, the investigation of the evolution of the crystal structure during the dealloying.When the solid-solution alloys are dealloyed, the formed NP samples usually inherit the crystal structure and orientation of their parent alloys, for instance, in Au-Ag alloys. 20,21This is expected from the dissolution/surface diffusion model since the singlephase solubility allows the composition to change continuously within the same crystal structure.In contrast, the dealloying of stoichiometric intermetallic compounds and amorphous alloys often induces the change in the crystal structure.][27][28][29][30][31][32] In this Letter, we investigated the evolution of the crystal structure during the dealloying of the Fe-Rh alloy, which exhibits a single bodycentered-cubic (bcc) phase with extended Rh solubility up to $50 at.% Rh. 33 However, we found that the bcc structure transforms to the facecentered-cubic (fcc) structure during the dealloying although the composition of the obtained NP sample still resides in the region of the bcc phase.The obtained NP Rh(Fe) samples have extremely small ligament sizes of only 2-3 nm and the retention of 60 at.% Fe, evidencing the stability of the NP structure against secondary coarsening.These findings were discussed within the framework of the dissolution/surface diffusion model, considered with the chemical disordering induced by the massive transfer of both elements during the dealloying.
We used the Fe-Rh alloy with a nominal composition of 85 Fe:15 Rh (at.%) as the parent alloy for the electrochemical dealloying.The alloy ingots were prepared by repetitive arc melting of Fe and Rh wires (>99.99%) in an Ar atmosphere.They were first homogenized at 1000 C for four days, compressed, and finally cut into millimetersized samples.Cold rolling was further used to prepare thin-sheet samples with a reduced thickness of $0.12 mm for in situ x-ray diffraction (XRD) measurements.All samples were annealed at 850 C to remove residual stress before the electrochemical dealloying.In both annealing steps, the samples were slowly cooled in the furnace that stabilized the samples in their equilibrium bcc phases.We used a three-electrode electrochemical system to characterize the anodization behavior and performed the electrochemical dealloying of the alloy samples in 1M H 2 SO 4 under potentiostatic control (Autolab PGSTAT302N).The working, counter, and reference electrodes were Fe 85 Rh 15 samples, Pt wires, and a pseudo Ag/AgCl electrode, respectively.The pseudo Ag/ AgCl reference electrode was prepared by anodizing an Ag wire upon application of 1 V against another Ag wire in 1 M HCl for 60 s.The potential of this peudo Ag/AgCl electrode was 0.420 6 0.002 V more positive than the standard Hg/HgO (6M KOH) electrode, and for comparison, all the voltages in the paper were converted to the Hg/ HgO scale.In situ x-ray diffraction was conducted to track the structural evolution with a parallel beam laboratory rotating anode diffractometer (Mo K a1,a2 radiation).The 2D diffraction images were integrated by pyFAI software and analyzed with the Rietveld method (TOPAS V6).NIST SRM660b LaB 6 powders were used for the detector calibration and the determination of the instrumental resolution.The microstructures of NP Fe(Rh) samples were characterized by transmission electron microscopy (TEM, FEI Titan 80-300) and scanning electron microscopy (SEM, Zeiss Ultra 600).When exposed to air, the as-prepared NP samples were very prone to oxidation due to the large amount of chemically active Fe elements and the nanometerscale ligaments.To avoid the oxidation during preparation of a TEM specimen, we ground the NP samples to fine powders inside a glovebox under argon protection, from where a vacuum-transfer TEM holder, containing the ground powder, can then be directly transferred to the TEM chamber for further observation.The residual Fe contents were measured by energy dispersive spectroscopy on the cross section of NP samples in a SEM chamber.
We first studied the electrochemical behavior of the Fe 85 Rh 15 alloy in 1 M H 2 SO 4 .As shown in Fig. 1, the anodic polarization curve reveals the existence of the critical dealloying potential (E critical ) and the passivation potential (E passivation ).As indicated by the arrow, E critical is at around À0.02 V, above which the current starts to increase exponentially, indicating the onset of the bulk dealloying.The existence of the critical dealloying potential has been observed in other binary solid-solution alloy systems such as the Ag-Au alloy. 34,35A series of chronoamperometry curves measured at different constant potentials around -0.02 V, which follows the procedures in Ref. 36, show that the dealloying can proceed at the lowest potential of À0.020 V.However, unlike in the case of the Au-Ag alloy, the dealloying current drops drastically at 0.23 V, referred to as the passivation potential.This arises from the formation of Rh oxide that prevents the exposure of the pristine alloy to the electrolyte, a phenomenon also observed in the dealloying of the Cu-Rh alloy. 37Consequently, the potentials from À0.020 V to about 0.23 V can be used to perform the electrochemical dealloying of Fe 85 Rh 15 alloys.Based on our previous study, 19 we chose the dealloying potential that is slightly higher than E critical , here À0.015 V, for the electrochemical dealloying of Fe 85 Rh 15 alloys.It minimized the influence of the secondary dealloying process and preserved the initial NP structure.
Figure 2(a) shows the evolution of the crystal structure of the Fe 85 Rh 15 alloy during the dealloying at À0.015 V by in situ XRD in transmission geometry.The transmission geometry allows the detection of the entire volume of the sample rather than only its surface (thickness $0.1 mm).At the beginning, the XRD pattern shows the typical bcc structure at room temperature.As the dealloying was commenced by applying À0.015 V, the intensity of bcc diffraction peaks decreased gradually and disappeared after $5 h.Meanwhile, another set of diffraction peaks appeared, which can be ascribed to an fcc structure.After about 10 h, the intensity of the newly-developed fcc pattern remained stable, indicating the completeness of the dealloying process.Figure 2(b) shows the diffraction patterns at three different stages of the dealloying process, i.e., the pristine alloy, the partially dealloyed sample, and the sample after the complete dealloying.It is clear that 1.An anodic polarization curve of the Fe 85 Rh 15 alloy in 1 M H 2 SO 4 , showing the existence of the critical dealloying potential and the passivation potential.E critical refers to the critical dealloying potential at À0.020 V, above which the current increases exponentially and the bulk dealloying can proceed.E passivation refers to the potential of 0.235 V, where the current drops drastically due to the passivation with Rh oxide.The scanning rate is 2 mV/s.after the complete dealloying, the diffraction peaks exhibit significant broadening, indicating the extremely small coherent diffraction domains, most likely the ligaments, and the presence of large microstrains.The contribution of the apparent grain size (D) and the micros-train (e) to peak broadening can usually be decoupled by Williamson-Hall analysis, i.e., b cosh ¼ Kk=D þ 4e sinh, in which b, K, and k are the integral width, Scherrer constant, and x-ray wavelength, respectively.Yet, a striking feature is the deviation of the Bragg peak positions from the ideal fcc structure, i.e., the 200 fcc peak is shifted to lower diffraction angles.These features indicate the existence of the anisotropic strain and/or stacking faults. 38The observed diffraction pattern resembles that of an expanded austenite, 39 whose crystal structure can be approximated with a faulted fcc structure and anisotropic distortions. 40To account for both the relative displacement between 111 fcc and 200 fcc and the peak broadening, we used an elastic anisotropy factor for cubic materials, which describes the anisotropic distortion 41 and the anisotropic peak broadening model with fitting parameters of S 400 and S 220 . 42Rietveld analysis shows that the ligament sizes are only about 3 nm.The refined anisotropic strain coefficients S 400 and S 220 indicate that the variance of the lattice parameters is largest in the [100] directions and smallest in the [111] directions with maximum micro-strain values of 0.8% and 0.3%, respectively.The micro-strain is much larger than that observed in NP Au with ligament sizes of $20 nm, which can result from the extremely small ligament sizes, due to which surface stress generates large stress and strain in the ligaments. 43It is worthwhile to note that the large micro-strain with anisotropic distortion leads to the co-existence of spin glass and ferromagnetic ground states in NP FeRh. 44igure 2(c) shows the change of the weight fractions of the bcc and fcc phases during the dealloying, which were obtained from the Rietveld fitting of the diffraction patterns during the dealloying [Fig.2(b)].Obviously, the fraction of the original bcc phase decreases linearly with the dealloying time, and correspondingly, the fraction of the fcc phase changes reversely.This linear change of phase fraction arises from the thin-sheet form of the Fe 85 Rh 15 alloy as shown in the inset of Fig. 2(c) due to which the dealloying proceeds along the out-of-plane direction with a constant rate.Our in situ XRD results, thus, clearly reveal the phase transformation of the Fe 85 Rh 15 alloy from the bcc to fcc structure during the dealloying.
We performed TEM to examine the microstructure of the NP Rh(Fe) samples.Figure 3(a) shows the typical NP morphology with interpenetrating ligaments and pores.The enlarged TEM images reveal the extremely small ligament sizes of only 2-3 nm, which evidences the suppression of the secondary coarsening during the dealloying [Fig.3(b)].The stabilization of the NP structure against coarsening results from the low surface diffusivity of Rh atoms, which plays the similar role as that of Pt atoms in Au(Pt)-Ag alloys. 19,45The corresponding selected area diffraction shows continuous diffraction rings, indicating the nearly randomly distributed nano-grains with the fcc structure [the inset in Fig. 3(b)].Moreover, energy dispersive spectroscopy shows that the Fe concentration increases monotonically with the lower dealloying potentials and reaches up to $60 at.% Fe at À0.015 V [Fig.3(c)].The retention of the high concentration of Fe elements is similar to that observed in the Pt-doped NP Au samples, in which 65 at.% Ag was observed when the secondary coarsening/dealloying was suppressed as well. 19We emphasize that the extremely small ligament sizes and the high concentration of residual Fe elements signify the preservation of the initial NP structure created by dealloying.Hence, the observed bcc-fcc transformation originates from the primary dealloying process rather than other secondary processes that occur behind the corrosion front.It is also worthwhile to note that the bcc-fcc phase transformation was observed in NP Fe(Rh) prepared by the dealloying at À0.005, 0.010, and 0.030 V, with Fe concentrations ranging from $50 at.% to 20 at.%.
Our in situ XRD and TEM results show the bcc-fcc transformation during the dealloying of the Fe 85 Rh 15 alloy.
Figure 4 shows that similar to Au-Ag alloys, Fe-Rh alloys have the single phase structure with the extended Rh solubility up to $50 at.%. 33 Since the composition of NP Fe(Rh) is $60 at.% Fe, which still resides in the bcc region, one may expect NP Fe(Rh) to retain the bcc structure of the parent alloy (arrow in Fig. 4).Instead, we observed the phase transformation from the bcc to fcc structure.In fact, as mentioned earlier, the resultant NP structure inherits the fcc structure of the parent alloys in the dealloying of Ag-Au alloys. 20,21owever, it should be noted that Au-Ag alloys show the single-phase solubility with the fcc structure across all temperatures and compositions, allowing for the continuous change in composition within the same crystal structure.However, in the case of the Fe-Rh alloy, there exists the phase transformation from the bcc to the fcc structure at high temperatures, where chemical disordering of Fe and Rh atoms is introduced.The chemical order indicates the occupation of the bodycentered sites by one kind of element and the corner sites by another kind, while chemical disordering introduces anti-site defects.It has been observed that the bcc-fcc phase transformation in Fe-Rh alloys can be triggered by the chemical-disordering process, for instance, by intensive ball milling. 46,47This means the bcc structure is the thermodynamically stable phase but can be converted to the metastable fcc structure when chemical disorder is introduced.
The observed phase transformation can be interpreted within the framework of the dissolution/surface diffusion model.4][15] In our case, Fe atoms are selectively dissolved during the dealloying of the Fe 85 Rh 15 alloy, while the remaining Rh and residual Fe atoms will diffuse along the alloy/electrolyte interface.The dissolution and surface diffusion processes can be highly dynamic and non-equilibrium, which will cause the disordered arrangement of Fe and Rh atoms at the alloy/electrolyte interface.The newly formed Fe-Rh clusters with the disordered arrangement of Fe and Rh atoms, thus, tend to form a metastable fcc structure rather than the equilibrium bcc structure.Consequently, the observed bcc-fcc transformation is driven by the chemical disordering of Fe and Rh atoms, as introduced by the highly dynamic dissolution and diffusion process.It is worthwhile to consider that due to the extremely small ligaments, surface stress may deform the ligaments immediately after their formation and even induce phase transformation.In this regard, the fcc structure is equivalent to a bodycentered-tetragonal (bct) structure with c/a ¼ ͱ2.Therefore, along the Bain path from bcc to bct, a tetragonal distortion of 41.4% would be required, 48 which is unlikely to happen in the considered case.
In summary, we studied the evolution of the crystal structure during the dealloying of Fe-Rh alloys that exhibit single-phase solubility up to 50 at.% Rh.We found that the crystal structure transforms from the original bcc to the fcc one during the dealloying.The resultant NP samples have extremely small ligament sizes of only 2-3 nm and the retention of a high Fe concentration of 60 at.%, which evidence the suppression of secondary coarsening/dealloying processes.The dealloying-induced phase transformation can be attributed to the disordered arrangement of Fe and Rh atoms, which is induced by the highly dynamic dissolution and surface diffusion processes at the alloy-electrolyte interface.Our study can sharpen the understanding of the spontaneous evolution process of nanoporosity and contribute to the development of novel NP metals.X.Y. would like to thank the Alexander von Humboldt Foundation.We acknowledge support by the KIT-Publication Fund of the Karlsruhe Institute of Technology.

FIG. 4.
Phase diagram of the Fe-Rh alloy showing the single-phase solid solubility up to $50 at.% Rh within the bcc structure.After the dealloying, the resultant NP Fe(Rh) is expected to have a bcc structure when the Rh concentration is below 50 at.%.However, the bcc-fcc transformation was observed although the resultant NP Fe(Rh) still retained 60 at.% Fe.

FIG. 2 .
FIG. 2. (a) A contour plot of in situ x-ray diffraction patterns during the dealloying of the Fe 85 Rh 15 alloy at À0.015 V, showing the phase transformation from the original bcc to fcc structure.(b) Rietveld fitting of the diffraction patterns at different stages of the dealloying process, i.e., the pristine alloy, the partially dealloyed sample, and the completely dealloyed sample.Note the significant peak broadening due to the formation of the NP structure.(c) The linear change in the weight fraction of the fcc and bcc phases during the dealloying, verifying the constant dealloying rate.Inset shows the photo of the NP sample with a thin-sheet form.

FIG. 3 .
FIG. 3. (a) bright-field TEM image showing the typical NP morphology with bi-continuous ligaments and pores.(b) The enlarged view of the NP morphology exhibiting a mean ligament size of only 2-3 nm.(Inset) The corresponding selected area diffraction patterns confirming the fcc structure.(c) The dependence of Fe concentrations on the dealloying potentials (E dealloy ).