more GBs decreased as the grain size got smaller
because the dislocation density saturated when
the strain fields of the GB surface terminations
overlapped. This finding parallels a long-standing
observation in metallurgy, in which attempts to
strengthen materials by simply decreasing grain
size yields progressively diminishing returns (43).
The saturation observed here contrasts with the
linear GB-activity relationships we previously observed over a large GB density range for NP catalysts. GBs in NPs typically have much narrower
strain fields that consequently have little overlap,
even at high GB density. (44, 45)
The identification of GB-stabilized dislocations as structural elements that increase local
activity implies that procedures that increase dislocation density should improve the activity of
bulk samples. Mechanical treatments are frequently used in metallurgy to harden metals by
inducing the formation of dislocation clusters
and new GBs (46, 47). Thus, we investigated the
ability of mechanical treatments to restore CO2
reduction activity to electrodes whose GB density had been reduced by annealing. Two strain-hardening treatments were applied to an Ann.
500 sample: hammering by hand on a small
anvil, and a more intensive cold-rolling procedure
(48, 49). We chose mechanical treatment as a
method to alter grain structure to minimize electrode roughening.
SEM images (fig. S10) indicate that the me-
chanically treated samples exhibit no long-range
mesostructuring. Measurement of the electrodes’
electrochemical surface area indicates rough-
ness factors of ~1 and ~1.2 for the hammered and
cold-rolled sample, respectively (table S2). By
EBSD (fig. S11), a decrease in grain size was ob-
served with deformation, concomitant with
changes in surface orientation, while dislocation
clusters were observed to begin forming new GBs.
Pb UPD voltammograms (fig. S11) indicated that
the parent Ann. 500 and the hammered sample
consisted predominantly of Au(111) surface facets,
whereas the cold-rolled sample exhibited a high
proportion of Au(110), in addition to a minor
population of Au(111).
The increased defect density in the mechani-
cally treated samples translated into substantial
increases in CO2 reduction activity. At –0.4 V
versus RHE, jCO for the hammered foil was higher
by a factor of ~2.3 than the parent Ann. 500 foil.
The cold-rolled sample, which was even more
defect rich, showed a jCO higher by a factor of ~3.9
than Ann. 500 and was ~50% more active than
un-annealed Poly-Au. (fig. S11). In contrast, jH2
decreased by ~30 and ~40% for the cold-worked
and hammered samples, respectively. Given the
large differences in faceting observed between
the hammered and cold-rolled samples, these
results demonstrated that the defects introduced
by mechanical treatment were responsible for
selectively increasing CO2 reduction activity. De-
spite possessing no mesostructuring to increase
local pH, or hydrophobic backing to enhance
CO2 transport, the jCO of cold-rolled sample was
within a factor of 2 of the highest surface area–
normalized activity that has been previously re-
ported for a nanostructured Au electrode (50).
We envision two possible explanations for the
increase in CO2 reduction activity at regions with
high dislocation densities. Lattice strain at the
surface induced by dislocations could alter the
binding energies for CO2 reduction intermedi-
ates in a way that reduces the overall barrier
(51, 52). Alternatively, dislocation surface termina-
tions may create high step densities that are more
active than terraces (53, 54). Resolving these and
other possible contributions will require atomic-
level structural information under operating
conditions. Although there are many possible
strategies for creating strained or stepped surfaces,
surface restructuring during catalysis can cause
rapid relaxation of high-energy surfaces. High-
energy surface structures that arise from bulk
defects are likely to be more resistant to such deg-
radation. The dislocation density of a polycrystal-
line material is determined not just by the GB
density but by the GB character distribution as
well (1, 55). Our results therefore motivate the
use of GB engineering to control these prop-
erties in heterogeneous catalysts.
REFERENCES AND NOTES
1. T. Watanabe, J. Mater. Sci. 46, 4095–4115 (2011).
2. A. Rollett, G. Gottstein, L. Shvindlerman, D. Molodov, Z. Metallk.
95, 226–229 (2004).
3. Y. Chen, C. W. Li, M. W. Kanan, J. Am. Chem. Soc. 134,
4. C. W. Li, J. Ciston, M. W. Kanan, Nature 508, 504–507
5. A. Verdaguer-Casadevall et al., J. Am. Chem. Soc. 137,
6. X. Feng, K. Jiang, S. Fan, M. W. Kanan, J. Am. Chem. Soc. 137,
7. X. Feng, K. Jiang, S. Fan, M. W. Kanan, ACS Cent. Sci. 2,
8. S. Choi et al., Chem. Cat. Chem. 7, 2077–2084 (2015).
9. X. Sun, K. Jiang, N. Zhang, S. Guo, X. Huang, ACS Nano 9,
10. X. Huang et al., Nano Lett. 14, 3887–3894 (2014).
11. C. Wang et al., Nano Lett. 16, 5669–5674 (2016).
12. B. D. Aaronson et al., J. Am. Chem. Soc. 135, 3873–3880
13. J. C. Byers, A. G. Güell, P. R. Unwin, J. Am. Chem. Soc. 136,
14. N. Ebejer et al., Annu. Rev. Anal. Chem. 6, 329–351 (2013).
15. A. G. Güell et al., Nano Lett. 14, 220–224 (2014).
16. D. Dingley, V. Randle, J. Mater. Sci. 27, 4545–4566 (1992).
17. F. Humphreys, J. Mater. Sci. 36, 3833–3854 (2001).
Fig. 4. HR-EBSD microstrain profiles of non-CSL and S3 GBs probed
in Fig. 3. Each color denotes a microstrain profile for a single grain, each
measured relative to a reference pixel at positions = 0 mm (for left, blue
grain) and 16.8 mm (red, right grain). (A to D) Left to right: lattice
microstrain profiles showing lattice distortion in the x, y, and z axes and
the EBSD orientation map showing the strain map region for the 17° GB
interrogated in Fig. 3, B to E. A large inflection is observed at ~8 mm,
the location of the GB. (E to H) Left to right: lattice microstrain profiles in
the x, y, and z axes and EBSD orientation map showing the strain map
for the S3 GB interrogated in Fig. 3, G and H.
RESEARCH | REPORTS