Toughening Elastomers with
Sacrificial Bonds and Watching
Them Break
Etienne Ducrot,
1,2,3
Yulan Chen,
4
Markus Bulters,
5
Rint P. Sijbesma,
4
Costantino Creton
1,2,3
*
Elastomers are widely used because of their large-strain reversible deformability. Most unfilled
elastomers suffer from a poor mechanical strength, which limits their use. Using sacrificial bonds,
we show how brittle, unfilled elastomers can be strongly reinforced in stiffness and toughness
(up to 4 megapascals and 9 kilojoules per square meter) by introducing a variable proportion of
isotropically prestretched chains that can break and dissipate energy before the material fails.
Chemoluminescent cross-linking molecules, which emit light as they break, map in real time where
and when many of these internal bonds break ahead of a propagating crack. The simple
methodology that we use to introduce sacrificial bonds, combined with the mapping of where
bonds break, has the potential to stimulate the development of new classes of unfilled tough
elastomers and better molecular models of the fracture of soft materials.
E
lastomers are widely used in industrial
applications such as tires, seals, gloves,
and dampers for their ability to deform
reversibly to large strains. Yet currently, main-
taining this lar ge deformability imposes an up-
per limit in stiffness. Above a Young’s modulus
of around 1 to 1.5 MPa, unfilled elastomers are
brittle, limiting applications. This limitation is
particularly severe at high temperatures and for
rubbers with a low entanglement density. Frac-
ture of simple elastomers has been described by
Lake and Thomas (1, 2), who predicted that the
threshold fracture toughness (minimum energy
necessary to break the elastomer) should scale
with N
c
1/2
,whereN
c
is the number of monomers
between cross-links. In essence, the more cross-
linked (and the stiffer) the elastomer, the more
brittle it becomes. T o circumvent this limitation,
rubbers have been toughened by adding nano-
fillers and by making the elastomers more visco-
elastic (3). It was found that optimized nanofillers
impart a large increase in stiffness at small strain
and cause highly dissipative processes to be active
at large strains, effectively increasing both stiff-
ness and strain at break (4). However, incorporat-
ing nanofillers introduces constraints in terms of
processing and safety and requires careful disper-
sion. Alternatively, increasing the viscoelastic char-
acter of the elastomer is used to increase fracture
toughness through molecular friction (5), but this
method only works over a limited temperature range.
Many empirical strategies have been tried to
increase simultaneously the strength and stiffness
of unfilled elastomers by introducing hetero-
geneities in the network. The multimodal dis-
tribution of molecular weights between cross-links
has been extensively tried on silicones, but the
gains in stiffness at low strains only result in
moderate increases in toughness (6–8). Another
strategy has been to prestretch a partially cross-
linked elastomer and further cross-link it in the
stretched state (9). This improves the strength of
the elastomer in the prestretching direction but
decreases the initial tensile modulus and leads
to very anisotropic properties. Knowledge-based
strategies to increase fracture toughness while
retaining full reversibility of deformation and
low viscoelasticity have been limited by the lack
of understanding of where and how e nergy is dis-
sipated as a crack propagates. This microm e c h a n -
ical knowledge has been developed decades ag o
for materials with a yield stress, such as metals
(10) and glassy polymers (11, 12), in which dis-
sipation of energy is localized and visible in a
plastically deformed zone. However, for soft ma-
terials (gels, rubbers, or soft adhesives) there is
typically no well-defined yield stress, and the
origin of dissipation of energy during crack
propagation remains an open question. Recent
work has shown hydrogels with toughness in
excess of 100 to 1000 times that of regular (very
brittle) gels while maintai ning a reasonable ex-
tensibility and recoverability of the strain (13, 14).
Hydrogels are very soft [Young’s modulus (E)
1 to 100 kPa] molecular sponges full of water
and thus are quite different materials from the
hydrophobic and fully water-insoluble elastomers.
The increase in toughness has been attributed
there to the early fracture of “weak or overstressed”
bon ds (either intrinsically weaker than the main
bonds, or loaded more than the main bonds) in-
troduced in the bulk of the material by means
of the synthetic method (13, 15–17).
Our elastomers are obtained through sequen-
tial free-radical polymerizations. First, a well–
cross-linked rubbery network is synthesized by
means of ultraviolet (UV) polymerization in the
presence of solvent. After drying, the rubber
sheet referred to as “single network” (SN) is then
swollen with a second monomer, UV initiator,
and a small amount of cross-linker, effectively
isotropically stretching the chains of the network.
A second UV polymerization is performed on
this swollen network, until all monomer is con-
sumed. Varying the degree of cross-linking of
the first network controls the level of swelling
from a swelling ratio (Q)=3.3to5andeffec-
ti ve ly co ntrols the level of prestretch of the first
network chains and their volume fraction in the
network (fig. S1). The double networks will
be referred to as “DN.” To obtain even lower-
volume fractions of first network and higher
levels of prestretch, the second step can be re-
peated on the DN. The presence of chains of the
second network decrease the elastic component
of the free energy per unit volume, allowing us
to stretch first network chains further. The third
polymerization step results in a material in
which the first network chains only represent
5 to 10 weight percent (wt %) and are highly
stretched, whereas the chains polymerized during
the second step are lightly stretched, and those
polymerized during the third step are entangled,
lightly cross-linked, and loosely connected by
chain transfer reactions (fig. S2 and table S3).
These triple networks will be referred to as TN.
It should be noted that because of the nature
of the monomers used (acrylates), chain trans-
fer reactions to the polymer likely occur during
the second and third polymerization, effective-
ly loosely connecting the networks with each
other and affecting stress transfer between the
networks.
In our example, the first network is made
from ethyl acrylate and butanediol diacrylate
cross-linker , and the second and third steps have
1
École Supérieure de Physique et de Chimie Industrielles
de la Ville de Paris (ESPCI) ParisTech, UMR 7615, 10, Rue
Vauquelin, 75231 Paris Cédex 05, France.
2
CNRS, UMR 7615,
10, Rue Vauquelin, 75231 Paris Cédex 05, France.
3
Sorbonne-
Universit és, Université Pierre et Marie Curie (UPMC) Université
Paris 06, UMR 7615, 10, Rue Vauquelin, 75231 Paris Cédex
05, France.
4
Institute for Complex Molecular Systems, Eindhoven
University of Technology, Post Office Box 513, 5600 MB
Eindhoven, Netherlands.
5
DSM Ahead, Urmonderbaan 22,
6167 RD Geleen, Netherlands.
*Corresponding author. E-mail: costantino.creton@espci.fr
been carried out with methyl acrylate. However,
similar networks have been obtained with se-
quential polymerizations by using two methyl
acrylate networks or two and three ethyl acrylate
networks with similar results (fig. S3 and table
S1). A summary of the properties and nomencla-
ture of the materials is shown in Table 1. Three
cross-linker concentrations in the first network
are presented here (table S1). EA
0.5
is cross-
linked at 1.45 mole percent (mol %) of mono-
mer , EA
1
two times more (2.81 mol %), and EA
2
four times more (5.81 mol %). Their correspond-
ing DN and TN are referred to as EA
x
MA if the
second network is made of MA or EA
x
EA if
the second monomer is EA, and referred to as
EA
x
MAMA if the third network is made of
MA or EA
x
EAEA if the third monomer is EA.
As an example of the mechanical properties
of SN, DN, and TN elastomers, the stress strain
curves at 60°C [45°C above the glass transition
temperature (T
g
) of the PMA] in uniaxial ext e n -
sion are shown in Fig. 1A for EA
1
,EA
1
MA, and
EA
1
MAMA. The elastic modulus of the elasto-
mer increases by a factor of up to 3.5, whereas the
true stress at break increases by a factor of 58.
The effect of changing the cross-linker concen-
tration in the first network on the properties of the
DN elastomer is shown in Fig. 1B, and the ef f ect
of changing the second and third monomer for the
DN and TN materials [at equal temperature dif-
ference from Tg. (T − T
g
)] is shown in Fig. 1C.
In principle, such a striking increase in strength
could be due to plasticity in large strain, leading
to permanent deformation upon unloading.
However, as shown in Fig. 2A the material
unloads with no detectable hysteresis for the
EA
1
MA DN, and as shown in Fig. 2B, there is
a substantial hysteresis during the first cycle but
no measurable hysteresis for the subsequent cy-
cles to the same strain for the EA
0.5
MAMA TN.
The residual deformation after each cycle remains
below 6% for the DN and TN, and the modulus
after each cycle remains nearly constant for the
TN, showing that the damage is very moderate
(fig. S4). Although substantial hysteresis occurs
in the material in cyclic extension, the modulus
does not decrease much. This suggests that unlike
what was observed for gels, most of the initial
modulus is due to the second and third network,
and the chains of the first network simply limit
the maximum extensibility.
The increase in strain and stress at break sug-
gests that the fracture toughness of the elas-
tomers should also increase. We carried out crack
propagation experiments on single-edged notched
samples (fig. S5) at a low nominal strain rate
of 0.025 s
−1
at 60°C and 20°C for all prepared
networks and used the large-strain approxima-
tion of Greensmith (18) to extract the critical
energy release rate G
c
. The results are presented
in Fig. 3 as a function of the elastic modulus. The
fracture toughness G
c
increases from ~50 J/m
2
to 2000 to 5000 J/m
2
from the SN and the TN.
These values of fracture toughness for materials
that also have a high elastic modulus and low
loading rate is in the range of some filled elas-
tomers or of the best tough hydrogels. W e are
aware of one unfilled elastomer that has a com-
parable combination of properties—namely, natural
rubber, which reaches a G
c
value of 2 to 10 kJ/m
2
in comparable conditions (19). But , this behavior
is due to strain-induced crystallization, which is
difficult to reproduce in other elastomers.
We hypothesize that the origin of the tough-
ening mechanism is similar to that of Gong and
coworkers (15, 20, 21)—that the fracture of
covalent bonds in the primary-minority network
controls the stress level (and hence the stiffness),
whereas the second- and third-majority networks
prevent large cracks from forming. From macro-
scopic evidence, Gong and coworkers could as-
sess that bonds actually break near the crack tip.
We found that we can dir ectly see where
and when sacrificial bonds break as the ma-
terial is deformed by using a chemoluminescent
Fig. 1. Mechanical behavior of tough elasto-
mers in uniaxial tension. (A) True stress/stretch
curves of the EA
1
,EA
1
MA, EA
1
MAMA, and PMA
second network alone at 60°C. (B) Effect of chang-
ing the level of cross-linker in the EA first net-
work on the EA
x
MA DN. (C) Comparisons between
EA
0.5
MA[MA] and EA
0.5
EA[EA] at equivalent T − T
g
.
Solid lines represent the TN, and dashed lines
represent the DN.
Fig. 2. Step-cycle loading-unloading curves of multiple network elastomers at 60°C. (A) Stress-
strain curve of a single sample of EAMA elastomer submitted to a step-cycle loading. (Inset) The applied
stretch as a function of time. All curves follow the same path on the s-l curve for this elastomer. (B)
ThesamegraphforasinglesampleofEA
0.5
MAMA network. In this case, each nth cycle follows a
different path when l of the nth cycle exceeds the maximum value of l of the (n – 1) cycle. Despite the
damage in large strain, the initial modulus is nearly the same for all cycles (fig. S4).
cross-linker, bis(adamantyl)-1,2-dioxetane bisacrylate
(BADOBA) (fig. S6) (22),whichisabletoemit
light when it breaks. If a sufficient force is ap-
plied to the BADOBA, the dioxetane group breaks
into two ketones, one of which is in the excited
state, as shown schematically in Fig. 4A. Relaxa-
tion to the ground state emits a photon in the bright
blue range of the spectrum (emission maximu m
Λ
max
= 420 nm). Because MA-based networks
need to be tested at 60°C and the BADOBA is
insufficiently stable at 60°C for the duration of
a mechanical test, we chose fully EA-based net-
works and prepared a SN of EA
0.5
,aDNof
EA
0.5
EA, and a TN of EA
0.5
EAEA. In all of
these networks, the first network was cross-
linked with BADOBA. Uniaxial tensile cycles
and fracture tests were filmed with a sensitive
camera able to detect single photons at 50 im-
ages per second (fig. S7). The load-unload cycles
of an EA
0.5
EA DN is shown in Fig. 4B, to-
gether with the corresponding light emission sig-
nal as a function of l,andthesamedataforthe
TN network is shown in Fig. 4, C and D. A
video of the luminescence of TN under cyclic
extension is presented in movie S1. Bond break-
age only occurs above a certain value of l,and
only for the first cycle, because subsequent cycles
to the same value of l are fully elastic. This data
demonstrate that the first-cycle hysteresis (fig. S8)
is due to the irreversible breakage of bonds ho-
mogeneously in the whole sample. Comparing
the mechanical data of Fig. 4, B and C, with that
of Fig. 1C, the replacement of the BDA with the
BA D O BA only has a sma l l wea k ening effect on
the mechanical properties of the elastomers, strong-
ly suggesting that although the chemoluminescent
bond is a bit weaker than a C-C bond (150 kJ/mol
versus 350 kJ/mol) (23), it really acts as a marker
and does not modify the mechanism itself. The
power-law correlation between the cumulative
mechanical hysteresis and the total emitted light
for a given value of l isshowninFig.4E.
Having established that the intensity of blue
light is directly connected to the mechanical hys-
teresis, the next step was to use it to map bond
breakage during fracture experiments. The image
of the luminescence around the tip of a propagat-
ingcrackisshowninFig.5fortheEA-basedSN,
DN, and TN. A video of the crack propagation
in the TN followed by chemoluminescence is
presented in movie S2. Because scales are
identical and the signal is proportional to the
number of photons per pixel, the intensity can
be compared; the bond breakage is very localized
in front of the crack tip (practically one pixel)
for the SN, is still localized at the crack tip but
more intense for the DN, and extends over a
large region in the material for the TN. A pre-
cise integ ration of the luminescence over the
whole region is not possible because of the high
dynamic range, leading to a saturation of the
camera sensor . However, these results show that
the same mechanism detected by the hysteresis in
uniaxial tension is active at the crack tip. This re-
sult is in qualitative agreement with post mortem
observations on fractured gels (20)andwithtwo
models proposed respectively by T anaka (24)and
Brown (25), predicting that the dual population of
co-continuous networks creates a yielding mech-
anism and a damage zone where the yield stress
is controlled by the stress to break the first net-
work bonds, and the size of the zone is controlled
by the extensibility of the second network chains.
However, the intensity mapping provides much
more precise information. It shows that an increase
in degree of prestretching of the chains and a de-
crease in their volume fraction leads to a much
Fig. 4. Chemoluminescent molecules to detect bond breakage. (A)Schematicofthechemo-
luminescence process, bond breakage of the dioxetane cross-linker, and light emission. (B) Step-cycle
test for a single sample of the DN network showing mechanical stress-st rain curves (red) and light emission
curves measured with image analysis (blue) for the same sample. A slight amount of bond breakage is
observed at high strain. (C) Step-cycle test for a single sample of the TN network showing light emission
(in blue) and stress (in red). Light emission is zero during unloading and reloading, until l of the nth cycle
exceeds the maximum value of l of the (n – 1) cycle. (D) Image of the luminescence of TN under loading
at various steps of the experiment [reported in (C)]: (1) near the beginning of the test, (2) on the first
loading at l ~ 2.6, and (3) on reloading at l ~1.7.(E) Cumulative light emitted during the cycles as a
function of the mechanical hysteresis in TN. The mechanicalhysteresisisdefinedassumoftheintegrals
under load-unload cycles, and the total light is the sum of the emitted light for that level of mechanical
hysteresis. The total light varies as the mechanical hysteresis to the power 0.75.
larger dissipative volume ahead of the crack tip
and to a tougher material, therefore guiding mate-
rials design. It also shows the dynamic shape of
the damage zone, allowing a quantitative com-
parison with more advanced damage models (26).
A family of tough, stiff unfilled elastomers
with less than 6% of residual deformation after
strainsupto150%,andnegligible viscoelasticity ,
can be obtained from ordinarily brittle elastomers.
The toughening mechanism relies on the dissi-
pation of energy due to bond breakage of a var-
iable fraction of sacrificial prestretched chains
inside the material. By varying the volume frac-
tion, monomer type, and cross-linking level of
the prestretched chains, the properties of the ma-
terials can be tuned over a wide range, and design
can be guided by the concomitant use of chemo-
luminescent molecules to reveal where and when
bonds break during fracture. The methodology can
be used both to develop better models of frac-
ture of soft materials and to guide design of
other families of soft materials than polyacrylics,
which have ordinarily poor mechanical proper-
ties but much better resistance to temperature,
UV, or chemicals.
References and Notes
1. G.J.Lake,P.B.Lindley,J. Appl. Polym. Sci. 9, 1233–1251
(1965).
2. G. J. Lake, A. G. Thomas, Proc. R. Soc. Lond. A Math.
Phys. Sci. 300, 108–119 (1967).
3. A. N. Gent, Langmuir 12, 4492–4496 (1996).
4. G. Heinrich, M. Kluppel, T. A. Vilgis, Curr. Opin. Solid St. M.
6, 195–203 (2002).
5. B. N. J. Persson, O. Albohr, G. Heinrich, H. Ueba, J. Phys.
Condens. Matter 17, R1071–R1142 (2005).
6. B. D. Viers, J. E. Mark, J. Macromol. Sci. Part A Pure
Appl. Chem. 44, 131 –138 (2007).
7. G. D. Genesky, C. Cohen, Polymer (Guildf.) 51, 4152–4159
(2010).
8. G. D. Genesky, B. M. Aguilera-Mercado, D. M. Bhawe,
F. A. Escobedo, C. Cohen, Macromolecules 41, 8231–8241
(2008).
9. N. K. Singh, A. J. Lesser, Macromolecules 44, 1480–1490
(2011).
10. G. E. Dieter, Mechanical Metallurgy (McGraw Hill,
New York, ed. 2, 1976).
11. H. R. Brown, Macromolecules 24, 2752– 2756 (1991).
12. E. J. Kramer, L. L. Berger, Adv. Polym. Sci. 91,1–68 (1990).
13. J. P. Gong, Y. Katsuyama, T. Kurokawa, Y. Osada,
Adv. Mater. 15, 1155– 1158 (2003).
14. Y. Tanaka et al., J. Phys. Chem. B 109, 11559–11562 (2005).
15. R. E. Webber, C. Creton, H. R. Brown, J. P. Gong,
Macromolecules 40, 2919–2927 (2007).
16. J.-Y. Sun et al., Nature 489, 133–136 (2012).
17. Y. H. Na et al., Macromolecules 39, 4641–4645 (2006).
18. H. W. Greensmith, J. Appl. Polym. Sci. 7, 993–1002 (1963).
19. K. Sakulkaew, A. G. Thomas, J. J. C. Busfield, Polym. Test.
30, 163–172 (2011).
20. Q. M. Yu, Y. Tanaka, H. Furukawa, T. Kurokawa, J. P. Gong,
Macromolecules 42, 3852–3855 (2009).
21. T. Nakajima, T. Kurokawa, S. Ahmed, W.- Wu, J. P. Gong,
Soft Matter 9, 1955–1966 (2013).
22. Y. Chen et al., Nat. Chem. 4, 559–562 (2012).
23. P. Lechtken, Chem. Ber. 109, 2862 –2870 (1976).
24. Y. Tanaka, Europhys. Lett. 78, 56005 (2007).
25. H. R. Brown, Macromolecules 40, 3815–3818 (2007).
26. X. Wang, W. Hong, Soft Matter 7, 8576–8581 (2011).
Acknowledgments: We gratefully acknowledge the financial
support of DSM Ahead, Geleen, Netherlands and many helpful
discussions with the researchers of their Materials Science Center.
We also thank E.J. Kramer and H.R. Brown for their insightful
comments and critical reading of the manuscript, and D. Martina
for his friendly and efficient help in setting up the camera for
the chemoluminescence experiments. Last, we thank W. Fresquet
from Andor Technology for lending us the camera.
Supplementary Materials
www.sciencemag.org/content/344/6180/186/suppl/DC1
Materials and Methods
Supplementary Text
Figs. S1 to S7
Tables S1 to S3
Reference (27)
Movies S1 to S2
13 November 2013; accepted 6 March 2014
10.1126/science.1248494
Mitosis Inhibits DNA Double-Strand
Break Repair to Guard Against
Telomere Fusions
Alexandre Orthwein,
1
Amélie Fradet-Turcotte,
1
Sylvie M. Noordermeer,
1
Marella D. Canny,
1
Catherine M. Brun,
1
Jonathan Strecker,
1,2
Cristina Escribano-Diaz,
1
Daniel Durocher
1,2
*
Mitotic cells inactivate DNA double-strand break (DSB) repair, but the rationale behind this suppression
remains unknown. Here, we unravel how mitosis blocks DSB repair and determine the consequences of
repair reactivation. Mitotic kinases phosphorylate the E3 ubiquitin ligase RNF8 and the nonhomologous
end joining factor 53BP1 to inhibit their recruitment to DSB-flanking chromatin. Restoration of RNF8
and 53BP1 accumulation at mitotic DSB sites activates DNA repair but is, paradoxically, deleterious.
Aberrantly controlled mitotic DSB repair leads to Aurora B kinase–dependent sister telomere fusions
that produce dicentric chromosomes and aneuploidy, especially in the presence of exogenous genotoxic
stress. We conclude that the capacity of mitotic DSB repair to destabilize the genome explains the
necessity for its suppression during mitosis, principally due to the fusogenic potential of mitotic telomeres.
W
e hypothesized that the inactivation of
mitotic double-strand break (DSB) re-
pair (1–3) might be caused by the fail-
ure to recruit the DSB repair factors BRCA1 and
53BP1 to DNA damage sites during mitosis
(4–7). 53BP1 and BRCA1 promote DSB repair
by nonhomologous end joining (NHEJ) and ho-
mologous recombination, respectively (8). Both
proteins accumulate at DSB sites downstream of
a common pathway consisting of the ataxia telan-
giectasia mutated–dependent phosphorylation of
H2AX (forming g-H2AX) followed by MDC1,
RNF8, and RNF168 recruitment (8). RNF168
ubiquitylates H2A (9, 10), which triggers the re-
cruitment of 53BP1 (11) and also BRCA1 (12).
Mitosis severs this pathway upstream of RNF8
recruitment to DSB sites as the formation of
g-H2AX and MDC1 ionizing radiation (IR)–
induced foci, which denote accumulation at break
sites, are unaffected by mitotic entry (4).
To elucidate how mitosis blocks RNF8 re-
cruitment to DSB sites, we tested whether the
RNF8-MDC1 interaction (13–15)isdisabledin
mitosis (see supplementary materials and meth-
ods). W e observed that whereas RNF8 and MDC1
interact in asynchronously dividing cells after ir-
radiation, their interaction is suppressed during
M (mitotic) phase (Fig. 1A and fig. S1A; details
of all synchronization and treatments are depicted
in fig. S2). Because RNF8 recognizes redundant,
phosphorylated epitopes on MDC1, we assessed
whether mitosis inhibits the ability of RNF8 to
recognize phospho-MDC1. MDC1-derived phos-
phopeptides that encompass its Thr
752
(T752)
phosphorylation site (pT752) are unable to re-
trieve RNF8 from mitotic extracts, whereas they
readily retrieve RNF8 from extracts of asynchro-
nously dividing cells (Fig. 1B and fig. S1B). This
inhibition is due to cyclin-dependent kinase
1(CDK1)–dependent phosphorylation, because
pretreating mitotic extracts with PP1, a Ser/Thr
phosphatase, or treating cells with the CDK1 in-
hibitor RO-3306 before harvesting restored the
RNF8-pT752 interaction (Fig. 1B).
To test whether CDK1 could directly inhibit
the RNF8-MDC1 interaction, recombinant hu-
man and mouse RNF8 proteins fused to glutathione
S-transferase (GST) were subjected to phospho-
rylation by CDK1-cyclin B or CDK2-cyclin A
before pull-down assays with pT752. CDK1, but
not CDK2, could phosphorylate RNF8, which
suppressed the ability of RNF8 to bind to pT752
(fig. S1, C to E). The reconstitution of the CDK1-
dependent inhibition of the RNF8-pT752 inter-
action identified T198 as the main CDK1 site on
RNF8 (fig. S1F). An antibody against the phos-
phorylated T198 residue (RNF8 pT198) confirmed
a mitosis-specific phosphorylation of T198 (Fig.
1C and fig. S1G). Mutation of T198 to alanine
(yielding RNF8 T198A) rendered the RNF8-
pT752 interaction insensitive to CDK1 in pull-
down assays (fig. S1E). In contrast, the T198E
(E, Glu) mutation, which mimics T198 phospho-
rylation, constitutively inhibited RNF8 binding
to phosphopeptides (fig. S1E).
We tested whether the RNF8-T198A mutation
restored RNF8 recruitment to DSB sites in mitotic
U2OS cells. Reintroduction of wild-type (WT) RNF8
![Fig. 4. Chemoluminescent molecules to detect bond breakage. (A) Schematic of the chemoluminescence process, bond breakage of the dioxetane cross-linker, and light emission. (B) Step-cycle test for a single sample of the DN network showing mechanical stress-strain curves (red) and light emission curves measured with image analysis (blue) for the same sample. A slight amount of bond breakage is observed at high strain. (C) Step-cycle test for a single sample of the TN network showing light emission (in blue) and stress (in red). Light emission is zero during unloading and reloading, until l of the nth cycle exceeds the maximum value of l of the (n – 1) cycle. (D) Image of the luminescence of TN under loading at various steps of the experiment [reported in (C)]: (1) near the beginning of the test, (2) on the first loading at l ~ 2.6, and (3) on reloading at l ~ 1.7. (E) Cumulative light emitted during the cycles as a function of the mechanical hysteresis in TN. The mechanical hysteresis is defined as sum of the integrals under load-unload cycles, and the total light is the sum of the emitted light for that level of mechanical hysteresis. The total light varies as the mechanical hysteresis to the power 0.75.](/figures/fig-4-chemoluminescent-molecules-to-detect-bond-breakage-a-1pcckvkk.webp)