Towards observing interstitial hydrogen in solution in metal lattices

Materials

The ingot of CoCrNiFeMn high-entropy alloy was prepared by vacuum arc melting a mixture of pure elements with purity greater than 99.95 wt.% under an argon atmosphere after a vacuum of 3.3 Pa has been reached. The ingot was melted at least four times to ensure compositional homogeneity, followed by homogenization at 1,200 °C for 24 h, cold rolling to a 90% thickness reduction (from 10 mm to 1 mm) and annealing at 950 °C for 4 h.

The sample with a nominal chemical composition of (CoCrNi) 94 Al 3 Ti 3 (at.%) was synthesized through vacuum arc melting high-purity metals (purity ≥ 99.9 wt.%) under an argon atmosphere. To ensure compositional homogeneity and prevent element segregation, the sample was melted five times before being poured into a round ingot with a diameter of 50 mm. The as-cast sample was homogenized at 1,200 °C for 3 h and then cold-rolled along with its thickness to reduce the thickness by 50%. Two different microstructures were achieved by applying distinct heat treatments to the as-rolled samples: heating at 1,000 °C for 1 h, yielding the precipitate-free state (solution-treated), followed by 800 °C for 5 h, yielding the precipitated sample. All heat-treated samples were water-quenched.

The ingot of CoNiV alloy was synthesized by vacuum arc melting under an argon atmosphere (99.99% purity) after a pressure of 3.3 Pa. To ensure compositional homogeneity, the ingots were remelted five times. The ingots were then hot-forged at 1,200 °C into rectangular slabs with a thickness of 5 mm and a width of 20 mm, followed by homogenization at 1,200 °C for 3 h. Subsequently, the homogenized alloys were cold-rolled to a total thickness reduction of 50%. Finally, the cold-rolled sheets were subjected to solution annealing at 1,000 °C for 1 h (solution-treated).

The Inconel 718 nickel-based superalloy had a chemical composition of 54.72Ni–18.60Fe–18.70Cr–4.24Nb–2.62Mo–0.80Ti–0.27Al–0.05C wt.% (53.93Ni–19.26Fe–20.80Cr–2.64Nb–1.58Mo–0.97Ti–0.58Al–0.24C at.%). This alloy (diameter 15 mm, length 70 mm) was fabricated using a commercial selective laser melting system (EOS M 400) with an average powder size of 30 ± 11 µm, a scanning strategy of 67° rotation, a laser power of 340 W, scanning speed of 1,100 mm s−1 and layer thickness of approximately 50 µm. The printed samples were initially homogenized at 1,100 °C for 1 h and then water-quenched to room temperature. Subsequently, the samples were subjected to a two-step ageing treatment: first at 720 °C for 1 h, followed by furnace cooling to 620 °C, at which they were held for 8 h before final air cooling to room temperature. The sequential D0 22 (γ″-type) and L1 2 (γ′-type) precipitates are introduced during the 720 °C and 620 °C ageing treatments, respectively.

Electron microscopy

Samples for the EBSD analysis were initially polished using 5000-grit SiC sandpaper, subsequently polished with diamond paste for 3 min and finally electrolytically polished in a 1:9 solution of HClO 4 and C 2 H 5 OH at 243 K for 20 s under a DC voltage of 25 V. The pre-thinned TEM foils, with a thickness of 50–60 μm, were polished using double-jet electropolishing at 243 K in 10 volume percent of HClO 4 in methanol to prepare the TEM sample.

EBSD mapping used a Symmetry S2 detector (Oxford Instruments) and TEM imaging was conducted on a JEM-2100F (JEOL). High-resolution STEM imaging was performed using the FEI Themis Z microscope equipped with an EDS detector. Atomic-resolution chemical characterization with compositional mapping was performed using four in-column Super-X silicon drift detectors.

Atom probe tomography

The sample preparation began with matchstick bars with dimensions 1 × 1 × 13 mm3. The samples were initially polished using 800-grit sandpaper. Coarse electropolishing was then performed using a 25% perchloric acid solution in glacial acetic acid at voltages ranging from 15 to 30 V to polish the samples into needle-like specimens. Fine electropolishing was subsequently performed using a 2% perchloric acid solution in ethylene glycol butyl ether at voltages ranging from 10 to 25 V. This fine electropolishing procedure used a platinum wire ring cathode and was monitored under a 40× optical microscope. A key step involved forming a neck until the top portion detached, creating a sharp tip.

Before deuterium charging, FIB tip sharpening was performed on the electropolished tips to ensure that the samples satisfied APT requirements in tip sharpness. APT was then performed using a LEAP 4000X instrument (CAMECA), operating in voltage-pulsing mode with a 0.2% detection rate, 200-kHz pulse frequency, 20% pulse fraction and sample temperature of 50 K. Data processing was performed using AP Suite software (version 6.1, CAMECA). The image compression factors and k factors used for reconstruction were 1.3 and 3.3 for the CoCrNiFeMn alloy and the Inconel 718 nickel-based superalloy and 1.5 and 3.3 for the (CoCrNi) 94 Al 3 Ti 3 alloy, respectively. These reconstruction parameters were selected following standard APT reconstruction procedures and, for the precipitate-containing datasets, refined by comparison with precipitate morphologies independently observed by TEM. Peak ranging was performed by considering the full isotopic distributions of the major alloying elements and possible molecular ions; the bulk composition analysis is shown in Supplementary Tables 2–4. A peak-free region around 2.5 Da was defined as noise and used to estimate the local low-mass background adjacent to the 2-Da signal. Three representative mass spectra are shown in Supplementary Fig. 8 and the results for all eight successful cryo-APT datasets out of 68 trials are summarized in Supplementary Table 5. Because D+ and H 2 + overlap at 2 Da and cannot be resolved directly, this overlap was treated separately as a systematic uncertainty and assessed using uncharged reference specimens analysed under identical voltage-pulsing conditions. APT reconstruction sensitivity was further evaluated by reconstructing the dataset shown in Fig. 3 using an alternative image compression factor value of 1.3, which did not alter the observed preferential deuterium enrichment within the L1 2 precipitate (Supplementary Fig. 9).

Electrolytic deuterium charging

A charging solution of 0.1 M NaOH in D 2 O (Sigma-Aldrich) was used to create a deuterium-rich environment around the samples at the cathode in an electrolytic circuit, with a gold counter-electrode. This was carried out in a custom-designed cryo-transfer glovebox (M. BRAUN Inertgas-Systeme). Four sharp specimens were pre-mounted in a four-needle atom probe sample puck with polyether ether ketone insulation. The four tips protruded from the puck to the same height and were immersed into the charging solution at a depth that was barely visible and could only be identified by means of the capillary of the liquid solution. Deuterium charging was conducted at 1.6 V for 90 s powered by Keysight 2901A, followed by immediately immersing the whole puck in a LN 2 bath.

Cryogenic sample transfer and cryo-FIB for APT specimens

After deuterium charging and LN 2 quenching, the specimens for Figs. 3 and 4 were transferred from the glovebox to a Hydra plasma focused ion beam (PFIB; Thermo Fisher Scientific) with a cryo-stage using a cryogenic sample carrier (Ferrovac). The sharpening process was completed on a cryo-stage at approximately −190 °C. Two ion-beam conditions were used—30 kV–0.1 nA and 30 kV–30 pA—to remove ice and further refine and smooth the tips during the annular milling process. A final polishing step at 5 kV–50 pA was applied for 10 s to remove the possible ion-damaged layer. The APT specimens were then transferred using the cryogenic transfer suitcase from the PFIB system to the APT system for APT measurements.

Thermal desorption spectroscopy

For the TDS tests, samples measuring 1 × 8 × 10 mm3 were polished using 5000-grit SiC sandpaper. The samples were then charged for 72 h at a constant current density of 20 mA cm−2 in a solution containing 0.5 M H 2 SO 4 and 2 g l−1 CH 4 N 2 S. TDS measurements were subsequently performed using a mass spectrometer (JTF-20A), with heating rates of 100 °C h−1, 200 °C h−1 and 300 °C h−1.

Density functional theory calculations

All DFT calculations were performed using the projector augmented wave39 method, as implemented in the Vienna Ab initio Simulation Package (VASP)40. The exchange-correlation effects were selected as the generalized gradient approximation determined by the Perdew–Burke–Ernzerhof functional41. To simulate chemical disorder, the special quasirandom structure (SQS) approach42 was applied using the Alloy Theoretic Automated Toolkit (ATAT)43,44. A plane-wave cut-off of 450 eV was used and the Brillouin zones were sampled with a gamma-centred Monkhorst–Pack grid45. Next, the Methfessel–Paxton method46 was used with a smearing width of 0.1 eV. The supercell size and atomic positions were allowed to relax until the Hellmann–Feynman forces were less than 0.01 eV Å−1 and convergence criteria for energies and forces were set to 1.0 × 10−5 eV per atom and 0.02 eV Å−1, respectively.

CoCrNiFeMn MPEA, CoCrNi MPEA and L1 2 Ni 3 (Al, Ti) precipitates, pure Ni, Ni 2 FeCr (matrix of Inconel 718 superalloy) and D0 22 Ni 3 Nb precipitates were studied in this work. The CoCrNiFeMn, CoCrNi, L1 2 Ni 3 (Al, Ti), Ni and Ni 2 FeCr are all FCC structures and D0 22 Ni 3 Nb has a tetragonal structure (I4/mmm, space group 139), the Ni atoms occupy the 4d (0, 1/2, 1/4) and 2b (0, 0, 1/2) positions and the Nb atoms occupy the 2a (0, 0, 0) positions. The SQS method was used to model the disordered phases including CoCrNiFeMn, CoCrNi, L1 2 Ni 3 (Al, Ti) and Ni 2 FeCr. The CoCrNi, Ni 2 FeCr, L1 2 Ni 3 (Al, Ti) and Ni phases were modelled using 108 atoms with 3 × 3 × 3 supercells according to their compositions, with all of the Ni atoms occupying the face-centred sites in L1 2 Ni 3 (Al, Ti). CoCrNiFeMn and D0 22 Ni 3 Nb were constructed by a 5 × 3 × 2 supercell containing 120 atoms and a 2 × 2 × 1 supercell containing 36 atoms, respectively. The CoCrNi (\(1\bar{1}\bar{1}\))/L1 2 Ni 3 (Al, Ti) (\(1\bar{1}\bar{1}\)) interface was constructed on the basis of experimental data and more than three SQS configurations for each bulk and interface structure were evaluated for comparison.

The DFT analysis mainly considered both octahedral interstitial and tetrahedral interstitial sites. The solution energy (E s ) of a hydrogen atom at a trapping site was calculated using equation (1):

$${E}_{{\rm{s}}}={E}_{{\rm{trap}}+{\rm{H}}}-{E}_{{\rm{trap}}}-1/2{E}_{{{\rm{H}}}_{2}}$$ (1)

in which E trap+H and E trap denote the total energies of the system with and without an interstitial hydrogen atom, respectively, and \(1/2{E}_{{{\rm{H}}}_{2}}\) denotes the half energy of a H 2 molecule. Given the heterogeneous chemical environment surrounding the interstitial sites in MPEAs, more than 20 randomly selected interstitial sites were sampled in each bulk phase structure, along with ten randomly selected interstitial sites at each interface structure, including both octahedral interstitial sites and tetrahedral interstitial sites.

We performed further calculations about the formation of vacancy. We selected three representative phases (CoCrNi, L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb) to conduct the relevant investigations, which can reflect the general energetic characteristics. Vacancies were introduced by removing atoms from the perfect supercell, the vacancy formation energy \({E}_{{\rm{V}}}^{{\rm{f}}}\) was computed as follows:

$${E}_{{\rm{V}}}^{{\rm{f}}}={E}_{{\rm{V}}}-{E}_{{\rm{bulk}}}+{\mu }_{i}$$ (2)

in which E V and E bulk are the energies of defective and defect-free supercells, respectively, and μ i is the chemical potential of the removed atom from the supercell. The chemical potential of the vacancy atom i was calculated according to the Widom-type substitution technique47. We calculated 40 substituted structures to determine the μ i of each element i. The calculated chemical potentials of different elements in CoCrNi were used to calculate the formation energies of vacancies in CoCrNi. For L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb, μ i denotes the total energy of the X atom at its ground state.

According to quantum mechanics, particles maintain a residual non-zero energy even at 0 K, which is defined as zero-point energy (ZPE)48. Here we used the ZPE and entropy corrections on H solution energies for three of the investigated alloys (CoCrNi, L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb). The Gibbs free energy of hydrogen dissolution, ΔG s is defined as follows:

$$\Delta {G}_{{\rm{s}}}={E}_{{\rm{s}}}+\Delta {E}_{{\rm{ZPE}}}-T\Delta S$$ (3)

in which E s is the calculated solution energy, ΔE ZPE is the difference in ZPE of hydrogen in the dissolved state and the gas phase and ΔS is the entropy difference between the dissolved state and the gas phase of hydrogen at standard conditions (atmospheric pressure, 298 K). The calculated ZPE of a diatomic H 2 gas is 0.269 eV, which is in good agreement with previous studies48,49,50. Supplementary Fig. 10 illustrates the comparison of energy results with and without ZPE and entropy corrections. The energy slightly increases after ZPE correction is applied and the increase of tet sites is marginally larger than the oct sites for CoCrNi and Ni 3 (Al, Ti). The energy increases further with more entropy corrections for all three alloys. It can be seen that the ZPE and entropy corrections exert no notable influence on the energy comparison.

Isotopic effect was also considered for these alloys (CoCrNi, L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb). H atom was substituted by D atom and the calculated results are shown in Supplementary Fig. 11. Replacing H with D exerts a negligible influence on the solution energies for all three phases. The energy slightly increases with the ZPE correction, which is similar to the case for H. However, the general trend of energy comparison is insensitive to isotopic effects. The diffusion of H/D atoms within the three phases (CoCrNi, L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb) was investigated using the climbing-image nudged elastic band method. As shown in Extended Data Fig. 10, the diffusion barriers of H and D in the three phases are similar to each other. After applying the corrections, the saddle points of the diffusion curves of all three phases (CoCrNi, L1 2 Ni 3 (Al, Ti) and D0 22 Ni 3 Nb) shift upward slightly, with minor differences observed between H and D. The general energetic trend remains the same after the corrections.

Quasi in situ TEM observation during electrochemical deuterium charging

Electropolished TEM foils of the aged (CoCrNi) 94 Al 3 Ti 3 alloy were used for quasi in situ D-charging experiments. Before charging, specific regions containing L1 2 precipitates were identified and characterized using bright-field TEM, high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM), selected-area electron diffraction (SAED) and HRTEM. The specimen was then electrochemically charged in 0.1 M NaOH prepared in D 2 O using a gold anode. A charging time of 1.5 min and the charging conditions used are the same as cryogenic APT experiments.

After charging, the specimen was removed from the electrolyte, dried and re-examined by TEM. The same regions and individual precipitates were relocated and compared with their former states. SAED and HRTEM analyses were performed after charging to assess whether D-charging introduced crystalline defects, including stacking faults, dislocations or other lattice defects.

Ion irradiation and positron annihilation spectroscopy

Ion irradiation was performed using a NEC 2 × 1.7 MV tandem accelerator at Peking University. The (CoCrNi) 94 Al 3 Ti 3 specimen (10 × 10 × 0.9 mm3) was irradiated with 3 MeV Ni2+ ions to a fluence of 1.5 × 1016 ions cm−2 at 500 °C, corresponding to a peak damage level of approximately 16 dpa, as calculated by SRIM-2013 using the quick Kinchin–Pease model. The displacement threshold energies used in the SRIM calculations were 40 eV for Co, Cr, Ni and Ti and 25 eV for Al. To minimize oxidation and other possible surface contamination during high-temperature irradiation, the target chamber was evacuated to a pressure below 1 × 10−4 Pa before heating and irradiation were initiated. After irradiation, the specimen was kept under vacuum during cooling and removed from the chamber only after the temperature had decreased to less than 50 °C.

Positron annihilation lifetime spectroscopy was performed using a 22Na positron source to complement the Doppler-broadening measurements and assess vacancy-type defects in the investigated alloys, corresponding to Supplementary Table 1. Vacancy-type defects introduced by hydrogen charging, performed under the same conditions as those used for the cryogenic APT experiments, and by ion irradiation were characterized by Doppler-broadening positron annihilation spectroscopy using a slow positron beam51,52 at the Institute of High Energy Physics, Chinese Academy of Sciences. A 22Na radioactive source was used to produce the positron beam, with a beam diameter of approximately 0.55 cm. To minimize specimen-to-specimen variability, Doppler-broadening positron annihilation spectroscopy measurements were performed sequentially on the same specimen in the pristine state, after electrochemical hydrogen charging and after subsequent Ni-ion irradiation. The incident positron energy varied from 0.5 to 25.0 keV. The corresponding mean positron implantation depth ranged from approximately 1.5 to 804.0 nm (corresponding to the irradiated region spanning approximately 5.8–15.3 dpa), as estimated using the empirical relation:

$$Z=\frac{40}{\rho }{E}^{1.6}$$ (4)