Accepted Manuscript

Ni-NiO composites obtained by controlled oxidation of green compacts

S. Cabanas-Polo, R. Bermejo, B. Ferrari, A.J. Sánchez-Herencia

PII: S0010-938X(11)00544-0

DOI: 10.1016/j.corsci.2011.10.016

Reference: CS 4664

To appear in: Corrosion Science

Received Date: 26 June 2011

Accepted Date: 12 October 2011

Please cite this article as: S. Cabanas-Polo, R. Bermejo, B. Ferrari, A.J. Sánchez-Herencia, Ni-NiO composites

obtained by controlled oxidation of green compacts, Corrosion Science (2011), doi: 10.1016/j.corsci.2011.10.016

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1

Ni-NiO composites obtained by controlled oxidation of green compacts

S. Cabanas-Poloa, R. Bermejob B Ferraria and A.J. Sánchez-Herenciaa,*

a.- Instituto de Cerámica y Vidrio (CSIC), C/ Kelsen, 5, 28049 Madrid

b.- Institut für Struktur- und Funktionskeramik, Montanuniversität Leoben, Peter-Tunner-

Straße 5, A-8700 Leoben, Austria

* Corresponding author: ajsanchez@icv.csic.es

ABSTRACT:

Ni-NiO composites have been obtained by the thermally induced oxidation of metallic

green compacts at temperatures between 300 and 450ºC and further sintering.

Thermogravimetric studies showed that oxidation process in air follows a quadratic

dependence with time for temperatures between 300 and 400ºC allowing the control in

the metal to oxide ratio. Microstructural analyses of compacts sintered in inert

atmosphere reveal a homogeneous distribution of phases. In the mechanical tests the

metal to ceramic ratio variation is evident in the ductile to brittle transition of the

fracture, making this method suitable to fabricate compacts with designed mechanical

properties.

Keywords: A. Nickel, A. Metal Matrix Composites, A. Ceramic Matrix Composites, B.

Weight loss, C. Oxidation, C. Oxide Coating.

1.- INTRODUCTION

The different mechanical behavior of metals and ceramics makes the CerMets a very

attractive way to overcome the structural problems that both materials have. In this

sense, the low toughness and mechanical reliability of ceramics can be improved by

the intimate incorporation of metals into the microstructure; while the low hardness and

wear resistance of metals can be addressed by the inclusion of ceramic particles in the

matrix [1, 2]. One of the main problems of the ductile and elastic balance in CerMets is

metal-ceramic bonding, which is usually low in the case of ceramic oxides. To solve

this problem several strategies have been followed aiming to design structures with

effective joining between both phases [3]. Approaches such as the employment of non

2

oxide ceramics which create compatible phases with the metal [4] or the use of metals

reinforced by their own oxides [5-8] have been proposed. The development of

interfaces mixing ceramics and the metal oxides has been another way to tailor

CerMets [9]. In the latter case the control of the oxygen plays an important role on the

development of the joining interface and consequently on the mechanical properties of

the composite [10]. The control in the partial oxygen pressure has been achieved either

by controlling the sintering atmosphere [11] or by sintering in a graphite bed which

allows setting the desired oxygen pressure.

Nickel-containing materials have a high technological interest in many aspects which

include construction and infrastructure, chemical industry, energy supply, or

transportation. These areas rely, to some degree, on nickel's unique combination of

properties. From a structural point of view, some of the most remarkable properties are

the high melting point (1453°C), the adherent oxide film formed on direct oxidation, the

resistance to corrosion by alkalis and its ductility conferred by the face-centered cubic

crystallographic structure. On the other hand, nickel oxide (NiO) presents unique

electrical, optical and magnetic properties. It is considered a prototypical p-type, wide

band gap (3.6–4.0 eV) semiconductor with special interests in applications such as

solar cells [12] or composite anodes for fuel cells, where certain mechanical properties

are required to confer structural integrity to functional devices [13]. Nevertheless, as a

ceramic material, the mechanical properties of nickel oxide are poor compared to some

oxide ceramics such as silica, alumina or zirconia. In contrast to the metallic nickel, the

nickel oxide presents higher hardness and lower fracture toughness [14, 15].

Therefore, Ni-NiO composites and structures with optimised mechanical properties are

pursued [5, 6].

Oxidation of nickel has been described for both compacts and powders in a wide range

of temperatures [16-21]. The growth of an oxide film on a metal by thermal oxidation is

usually discussed in terms of Wagner´s theory, in which the oxidation rate is controlled

by the transport of ions across the oxide film under the combined effect of

concentration gradients and electric fields. This leads to a parabolic kinetics with the

growth of the oxide layer. Oxidation of nickel is generally accepted to be governed by a

diffusion mechanism through the oxide layer formed on the nickel particle. The thicker

the oxide layer grows, the lower the oxidation rate becomes. Based on the assumption

3

of a diffusive process, long thermal treatments are needed to ensure a uniform

oxidation layer of the nickel.

In the present work the fabrication of Ni-NiO composites is approached by the direct

oxidation of green compacts of nickel with open porosity, and further sintering.

Oxidation has been induced by thermal treatments in air during twelve hours at

different temperatures, i.e. 250, 300, 350, 400 and 450 °C, in order to ensure a stable

and reliable metal to metal oxide ratio. Microstructural analysis of the different

specimens has also been considered to determine the distribution of metallic and

ceramic phases depending on the pre-oxidation temperature. Finally, the fracture

behavior of the samples has been assessed on notched specimens to evaluate the

influence of the Ni-phase on the crack growth resistance of the composite.

2. MATERIAL AND METHODS

Starting powders were commercial high purity nickel (INCO T-110, Canada) with a

mean particle size in volume of 2.5 µm and a specific surface area of 1.0 m2/g.

Concentrated aqueous suspensions of Ni powders were prepared in deionized water to

solids loadings of 40 vol.% (85.5 wt.%). A commercial polyelectrolyte (Duramax D-

3005, Rohm & Haas, USA, molecular weight of2,400) was used as dispersant. The

amount of dispersant employed was 1 wt.% on a dry solids basis and pH of the slurry

was adjusted to 10 by adding tetramethylammonium hydroxide. Green plates of 70 x

70 x 10 mm3 were fabricated by slip casting of the nickel aqueous slurries on a plaster

of Paris mould. A more detailed description of the processing can be found elsewhere

[22-23]. Density of the green compacts was determined by the Archimedes method in

mercury.

The thermal behavior of the nickel compacts was investigated on pieces cut from the

cast plates was determined using a thermo-gravimeter analyzer (ATD-TG, STA 409,

Netzsch, Germany) with a heating rate of 5 °C/min. Treatments of the nickel green

plates prior to sintering were made on a muffle furnace at 250, 300, 350, 400 and 450

°C with the same heating rate. In order to get a uniform oxidation of particles, thermal

treatments were carried out for 12h. The weight change has been calculated considering

the increment of weight (∆w = (w-w0)) and normalized by the starting weigh (w0) in each

case. Samples were weighted before and after the thermal pre-treatment employing an

4

analytical scale with a precision of 0.1 mg in order to calculate weight changes and the

correspondence with thermo-gravimetric data. Thermally pre-treated samples were

sintered in a tube furnace at 1460 °C under inert flowing Argon atmosphere. Density of

the sintered samples was measured by the Archimedes method in water. Sintered

samples were ground and polished with diamond paste of 1 µm for microstructural

characterisation. Specimen microstructures were examined by optical light reflected

microscopy (MOLR) in a DSM-950 microscope (Zeiss, Germany) and a field emission

scanning electron microscopy (FE-SEM) in a S-4700 microscope (Hitachi, Japan), In

order to avoid microstructural evolution, no further treatment was made on the samples

after the polishing. MOLR was performed on the polished surface of the sintered

samples. X-ray diffractograms (XRD) were recorded on the same polished surfaces

using a Siemens D-5000 diffractometer (Germany) that employs the Kα radiation of Cu.

For mechanical characterization (i.e. hardness and fracture toughness) prismatic

specimens were cut from the plates. Vickers hardness (Hv) was determined by

measuring the impression made with a Vickers indenter (Zwick 3212B, Germany) at an

applied load of 49 N and holding time of 10 s. A total of ten indentations were made on

each sample. The fracture toughness(KIC) of the different composites was determined

using the Single Edge V-Notch Beam method (SEVNB) on standard specimens (5

specimens from each class) of dimensions 3 x 4 x 45 mm3. A razor blade automatic

machine was utilized to create the notch, which was sharpened to a radius of less than

10 µm in order to minimize the influence of notch radius on KIC values [24]. The

notched specimens were then fractured under four-point bending under displacement

rate of 0.5 mm/min (environmental conditions of 25 °C and 40 % RH) on a standard

testing machine (Zwick Z010, Zwick/Roell, Ulm, Germany) according to ENV-843-1

[25].

3. RESULTS AND DISCUSSION

3.1. Oxidation of nickel compacts

Fig. 1 shows the FE-SEM micrograph taken on the fracture surface of the green

compact. In contrast to the green samples fabricated by pressing, the starting green

compacts fabricated by slip casting of aqueous slurries are composed of spherical

particles packet to a 55% of the density of the starting nickel powder (8.75 g/cm3

Determined by helium picnometry). It can be seen that the absence of pressure

5

avoided the deformation of the metallic particles and maintained the open porosity,

ensuring the gas diffusion toward the interior of the sample.

Fig. 2 shows the thermogravimetric curve of a pure nickel sample recorded in air from

room temperature up to 1400 ºC. At the beginning of the heating process, up to 300 ºC,

a weight loss that reaches 1 % of the original weight can be appreciated. It can be

attributed to the removal of the remaining water and carbon dioxide attached to the

surface [26, 27] as well as the dispersant employed to stabilize the aqueous nickel

slurries. As the temperature increases up to the range between 250-350 ºC (insert in

Fig. 2), a competition between this mass loss and the mass gain due to the oxidation of

the metallic particles is observed. Above this temperature the nickel oxidation prevails

and the weight rapidly increases until 650 ºC where it stabilizes. From the derivative

curve (see additional figures in the online edition) an inflexion point at nearly 400 ºC is

evidenced, where the oxidation rate slightly decreases. This point is easily overcome

and the maximum oxidation rate is achieved at 450 ºC. Finally, there is another mass

increment over 800 ºC which can be attributed to the final oxidation of large metal

particles. This late oxidation process is due to deviations on the diffusion kinetics at

high temperatures, also observed by other authors [28-30]. The total weight gain during

the oxidation is about 20.5 %, not very far from the theoretical value of 26.3 % for a

complete oxidation, but indicating that metallic nickel remains inside the sample. This

has been previously demonstrated by XRD for nickel compacts fabricated by other

colloidal techniques [31].

In order to achieve a better control of the oxidation of the particles that form the porous

green body, samples were isothermally treated in air at 250, 300, 350, 400 and 450 ºC

for 12 h. Fig. 3 shows the thermogravimetric curves for these treatments, including the

heating process. In all cases the heating step at 5 ºC/min shows a weight loss around 1

% in coincidence with that observed in fig. 2. As discussed before, it corresponds to the

removal of organic matter as well as absorbed water and gases. The dwell period at

250 ºC shows a very slightly weight variation after the aforementioned loss, keeping

the total weight unchanged after the whole cycle. In the case of the sample treated at

300 ºC after the start of the heating period, a light weight gain can be observed,

confirming that oxidation starts to be detectable between 250 and 300 ºC. For the

treatments made at higher temperatures it is clearly seen the weight gain due to the

oxidation of nickel and how the reaction grade increases with temperature, since it is a

diffusion controlled process [16, 18]. At 450 ºC the oxidation degree at the end of the

6

dwell period is close to those observed in the calcination process at high temperatures

described in fig. 2, also indicating that a metallic core remains.

To establish the kinetics of the oxidation process the thermogravimetric curves were

fitted to a parabolic equation, considering only the weight variation during the

isothermal stage, i.e. without taken into account the weight variations during the initial 5

°C/min heating stage. Fig. 4 shows the evolution of the square of the normalized weight

gain (∆w/w0) 2 over time and its mathematical fit to a parabolic equation for each

thermal treatment. It is observed that annealing at 250 ºC is far from a parabolic

behavior (fig 4.a). In fact, at this low temperature, the weight gain follows a linear fit

(R=0.997) with the time indicating that the oxidation process follows an interface-

controlled mechanism during this period. In the case of the samples treated at 300 and

350 ºC (fig. 4.a) weight gains clearly follow a parabolic behavior, which agrees with a

diffusive process of nickel along the grain boundaries in the NiO crust. For

temperatures of 400 and 450 ºC the oxidation shows deviations from a parabolic trend,

being very significant at 450 ºC (fig. 4.b). This behavior has been explained based on

three different factors related to the oxidation process and the particulate nature of the

green compacts. First at all, nickel particles join due to the oxidation of nickel which

inhibits the diffusion of the oxygen through the porosity of the sample [21, 32]. Then,

and later Ni vacancies diffuse along grain boundaries [33]. Finally, the residual

compressive stress, due to the expansion associated with the oxidation, governs the

behavior of small particles [18]. In fact, thermogravimetric studies of oxidation

performed on dense foils indicated that oxidation at this temperature is controlled by a

diffusion process with a parabolic kinetic [20].

The different Ni/NiO ratios achieved after the thermal pre-treatment for 12 h are

collected in Table 1. As expected, the final amount of oxide depends on the

temperature, allowing the fabrication of green compacts with compositions varying from

metal matrix composites to ceramic matrix composites.

3.2. Microstructural characterization

A set of prismatic samples pre-oxidized at 250, 300, 350, 400 and 450 ºC were

sintered at 1460 ºC under flowing argon atmosphere for 2 h. It should be noted that, at

7

the selected sintering temperature, the metallic nickel is melted and consequently

samples treated at 250 and 300 ºC loose their shape and a big metallic drop is

obtained after the thermal cycle. However, in samples pre-oxidized at higher

temperatures, the oxide layer formed on the particles surface creates a skeleton

structure that supports the original prismatic shape when nickel melts. As melted nickel

wets its own oxide, capillarity forces maintain the liquid within the oxide skeleton

avoiding mass loss of melt nickel. After sintering, the different Ni/NiO ratios are

maintained within the samples. This can be observed in Fig. 5, where the XRD patterns

recorded on cross-sections of pre-treated and sintered samples are presented.

Different ratios of Ni and NiO peaks are clearly observed.

Fig. 6 shows the microstructure of sintered and polished samples. The dark phase

corresponds to NiO, whereas the bright one refers to metallic Ni. It can be seen that

depending on the temperature of the pre-oxidation treatment, different microstructure in

terms of amount, size and distribution of NiO in the composite can be obtained. The

sample pre-oxidized at 300 ºC shows a very discrete NiO phase sparsely dispersed in

a continuous metallic phase. The particle size of the oxide grains is very uniform with

an average value of 5 ± 1 µm. As the temperature of the treatment increases to 350 °C,

the NiO phase grows within the Ni phase to form globular grains of 25 µm which start to

neck between them, forming a continuous ceramic phase of NiO. For this sample there

is still a continuous metallic phase, whereas for the sample at 400 ºC the ceramic

phase has become the continuous one. The presence of ceramic grains in the metallic

phase could be explained in terms of an oxide aggregation process taking place at the

same time as the thermal oxidation [34]. Joining of the oxide crust develops a structure

into the green compact, which is enlarged as the temperature of the thermal treatment

increases. As the sintering progress and the melting temperature of nickel is overcome,

the resulting structure of oxide aggregates and forms the rounded nickel oxide grains

observed inside the nickel matrix. For the sample at 450 ºC the NiO phase fills almost

the whole sample except for a few Ni grains that remain well dispersed within the

composite. The remaining Ni phase agrees with figs. 2 and 3 in the sense that the

oxidation rate does not reach the theoretical value for a complete oxidation. In this last

sample small holes are observed corresponding mainly to the pull out of grains during

the polishing process.

3.3. Fracture behavior

8

Mechanical characterization was only performed on samples pre-treated at 350 °C or

higher. The results obtained for Vickers indentations for hardness measurement in

composites pre-oxidized at 350, 400 and 450 ºC are summarized in Table 2. Hardness

data showed a small dispersion of data. Figure 7 shows the imprints corresponding to

the indentations in the three composites. It can be seen that the indentation covers a

large amount of grains of both Ni and NiO phases. As expected, the Vickers hardness

increases with the amount of ceramic phase (NiO), i.e. as the treatment temperature

increases, indicating a higher “ceramization” of the sample.

The fracture toughness evaluated following the VAMAS procedure on pre-notched

specimens was determined as [35]:

YW

SS

WB

FaYK io

Ic ⋅−

⋅−⋅=⋅⋅= 5.1)1(23

δδσ (1)

where σ is the stress, F the fracture load (in N), So and Si are the outer and inner spans

respectively (in m), B and W are the specimen thickness and width respectively (in m)

and Y is a geometric factor defined for an edge crack of length a as:

2

2

)1(

)1()35.168.049.3(326.19887.1

δδδδδδ

+−⋅⋅⋅+⋅−−⋅−=Y (2)

with Wa=δ

The results are also reported in Table 2. We caution the reader that the fracture

toughness measurements performed in the 350 °C pre-oxidized specimens are not

given in Table 2, since the SEVNB method on pre-notched specimens is no longer

valid due to the presence of plasticity. The determination of fracture toughness would

require a pre-cracked specimen (e.g. using the fatigue pre-cracking method), which is

out of the scope of this work.

9

The different fracture behavior of the three composite samples can be observed in the

load-displacement curves shown in Fig. 8. Whereas specimens pre-oxidized at 400 °C

and 450 °C show a typical brittle behavior (i.e. crack propagates unstable when the

applied stress intensity factor overcomes the fracture toughness of the material),

specimens pre-oxidized at 350 °C show a change in the compliance in the load-

displacement curve. It should be noticed that in this plot the slope corresponding to

samples treated at 400°C is similar to that of the samples treated at 450ºC, but with a

higher fracture load. This behavior indicates that in samples treated at 400°C, the

ceramic phase is continuous and the higher facture value is due to the higher metal

content. In fact, as it can be inferred from the microstructures (see Fig. 6), for such

composites the Ni forms a continuous phase (globular grains) and the NiO ceramic

phase is dispersed between the metallic grains. The macroscopic effect of the metal

phase on the crack propagation can be seen in Fig. 9. A clear difference can be

recognized between the three composites in terms of crack path. Whereas fractured

specimens pre-oxidized at 450 °C (Fig. 9c) show a straight crack propagation, a more

tortuous crack path can be identified in the ones pre-oxidized at 400 °C (Fig. 9b) and

specially at 350 °C (Fig. 9a). This can be explained by observing the fracture surfaces

of the three composites (Fig. 10). The 350°C specimen (Fig. 10a) shows the fracture in

the NiO phase, which is surrounded by a relative high content of Ni. Although the

globular NiO grains beak in a brittle manner, some ductile fracture can be infered in the

Ni phase. In the 400°C samples (Fig. 10b), the NiO grains are no longer globular but

rather equiaxial. The Ni phase results lower but is well distributed between the NiO

grains. Traces of brittle fracture can be clearly seen in the NiO grains. Finally, the

450°C samples show almost no remaining Ni phase (Fig. 10c). The fracture is very

brittle, as can be observed on the NiO grains.

In order to compare the three materials in terms of resistance to crack propagation, the

work of fracture (γWOF) was derived from the load-displacement curves showed in

Fig. 8. It is assumed that all the work done was consumed in the growth of the crack

and no significant elastic energy was stored in the specimen during testing. An

assessment of the total work of fracture (γWOF) is attempted through integration of the

registered load-displacement curve divided by twice the cross-section of the specimen

according to [36]. The corresponding results are given in Table 2. As expected, the

γWOF measured for the composite pre-oxidized at 350 °C is significantly higher than the

one determined for the specimens pre-oxidized at higher temperatures. Very

interesting is the difference between composites oxidized at 400 °C and 450 °C.

10

Whereas the latter shows a low γWOF, the former reveals a relative high level of γWOF,

while maintaining the stiffness associated with the NiO matrix.

Based on the fracture toughness values and γWOF, as well as on the crack path

observed on the fractured specimens, it can be concluded that specimens pre-oxidized

at 400 °C are good composite candidates to improve the mechanical properties (i.e.

fracture toughness) of NiO based ceramics through the presence of Ni metallic phase,

while maintaining their stiffness and hardness.

4. Conclusions

Ni-NiO composites pre-oxidized at different temperatures, i.e. from 250 to 450 °C, have

been investigated to optimize their microstructure-related mechanical properties. The

thermogravimetric curves of Ni have revealed that: (i) oxidation follows an interface-

controlled mechanism at low temperatures (i.e. 250 °C) that fits to a linear weight gain

with time, (ii) at higher temperatures the mechanism evolves to a diffusion-controlled

process making the final amount of Ni-NiO dependent on the oxidation temperature.

This final amount of NiO varies from 11 vol.% for 300 ºC up to 87 vol.% for 450 ºC. (iii)

The compacts do not achieve the complete oxidation, indicating that a nickel core

always remains within the oxidized particles.

A microstructural analysis of the different samples showed very different distribution of

metallic and ceramic phases. These different metal-oxide ratios are responsible for the

different fracture response of the materials. Whereas a metal-like behavior is achieved

for samples pre-oxidized at 350 °C and a very brittle fracture is exhibited by samples

pre-oxidized at 450 °C, an enhanced fracture toughness (and work of fracture) can be

achieved for 400 °C samples (i.e. KIC = 5.2 MPa·m1/2) with a relative high hardness of

2.0 GPa. This study shows that tailoring the microstructure of Ni by controlling pre-

oxidation conditions is possible, which allows optimizing the mechanical properties of

Ni-NiO composites for a specific application.

Acknowledgements

This work has been supported by MICINN-Spain under contracts IPT-310000-2010-12

and MAT2009-14448-C02-01.

11

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[33] A. Atkinson, R.I. Taylor, P.D. Goode, Transport Processes in the Oxidation of Ni Studied Using Tracers in Growing Nio Scales, Oxid. Met. 13 (1979) 519-543.

[35] ISO/DIS 23146 Fine ceramics (advanced ceramics, advanced technical ceramics) – Test methods for fracture toughness of monolithic ceramics – Single-edge V-notch beam (SEVNB) method. 2007

[36] H.G. Tattersall, G. Tappin, The work of fracture and its measurement in metals, ceramics and other materials, J. Mater. Sci. 1 (1966) 296-301.

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Caption to figures

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Fig. 1.- FE-SEM micrograph on the fracture surface of a green sample of pure nickel

fabricated by slip casting of metallic slurries .

Fig. 2.- Thermogravimetric curve of a Nickel green sample in air .

Fig. 3.- Weight change for different nickel compacts heated in air at a fixed

temperature for 12 hours.

Fig. 4.-. Evolution of the square of weight change over time for thermal treatments at

250, 300 and 350ºC (a) and 400 and 450ºC (b). Mathematical fits of experimental

results to a line have been also included.

Fig. 5.- X-Ray diffractograms recorded on the polished surface of pre-oxidized samples

sintered at 1460 °C in flowing argon.

Fig. 6.- Micrographs of the Ni-NiO microstructures of different samples oxidized at

300, 350, 400 and 450 °C.

Fig. 7.- Vickers indentation imprints (5 kg) on Ni-NiO specimens pre-oxidized at a) 350

°C, b) 400 °C and c) 450 °C.

Fig. 8.- Load - displacement curves of three composites pre-oxidized at different

temperatures (i.e. 350 °C, 400 °C and 450 °C).

Fig. 9.- SEVNB fractured specimens pre-oxidized at a) 350 °C, b) 400 °C and c) 450

°C.

Fig. 10- SEM micrographs of the fracture surface corresponding to SEVNB tested

specimens pre-oxidized at a) 350 °C, b) 400 °C and c) 450 °C.

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Table 1. Values of weight change and associated Nickel oxide amount of the samples

pre-treated at different temperatures.

Treatment Temperature [°C] 300 350 400 450

Weight Change [%] 1.8 5.2 13.8 21.6

NiO [vol. %] 10.7 28.6 63.5 86.6

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Table 2. Mechanical properties of three Ni-NiO composites pre-oxidized at 350°C,

400°C or 450°C respectively.

Sample pre-

oxidation

Vickers

Hardness

(HV5) [GPa]

Fracture Toughness

(KIc) [MPam1/2]

Work of fracture

(γWOF) [J/m2]

350 °C 1.1 ± 0.1 --- 117 ± 3

400 °C 1.95 ± 0.06 5.2 ± 0.8 55 ± 1

450 °C 3.03 ± 0.04 1.5 ± 0.1 10 ± 1

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Graphical Abstract