Hydration of alkali-activated slag: comparison

with ordinary Portland cement

A. Gruskovnjak,* B. Lothenbach,* L. Holzer,* R. Figi* and F. Winnefeld*

EMPA, Materials Science and Technology, Dubendorf, Switzerland

A multi-method approach was used for the investigation and comparison of alkali-activated slag binders (AAS),

pure slag and ordinary Portland cement (OPC). X-ray fluorescence, X-ray powder diffraction, granulometry,

calorimetry, thermo-gravimetric analysis and environmental scanning electron microscope investigations of the

microstructure with energy dispersive X-ray analyses were used to characterise the cements and their hydrate

phases. In addition, the chemical composition of the pore solution, including the different sulphur-containing ions,

was analysed. The precipitation mechanisms during binder hydration in the AAS and OPC systems exhibit

significant differences: in AAS the formation of the ‘outer product’ C–S–H is much faster than in OPC. The high

Si concentrations in the pore solution during the early hydration of AAS are related to the fast dissolution of

Na-metasilicate. The fast reaction of Na is an important factor for the voluminous precipitation of C–S–H within

the interstitial space already during the first 24 h. In addition to the Na-metasilicate component, the high fineness

of the slag represents a further important factor for the fast hydration of AAS. The small slag particles (, 2 �m)are completely dissolved or hydrated within the first 24 h, whereas hydration of the larger particles is much

slower. The fast formation of a gel-like matrix in AAS is the product of a fast ‘through solution’ precipitation,

which contrasts with the slower dissolution–precipitation mechanism of a ‘topotactic’ growth of C–S–H in OPC.

The chemical and mineralogical characterisation of solid and liquid phases and their changes with time are the

basis for thermodynamic modelling of the corresponding hydration process, which is presented in a second paper.

Introduction

Ground granulated blast-furnace slag can be acti-

vated with alkalis to obtain a clinker-free binder. The

production of alkali-activated slag (AAS) utilises in-

dustrial by-products, requires less energy than ordinary

Portland cement (OPC) and is associated with low CO2

emission.

The reactivity and strength development of slag

depend on various factors such as composition, glass

content, particle size distribution, type of activator and

alkali content.1,2 Different activators such as alkali

hydroxides, Na-metasilicate (waterglass) and others

have been used.3–6 Na-metasilicate was found to be the

most efficient2,7 activator which results in rapid hard-

ening and high compressive strength in comparison with

OPC.8 Slag cements activated with Na-metasilicate

[Na2O(n)SiO2] with moduli n between 0.6 and 1.5

showed higher ultimate strengths than Portland cement.9

The main hydration products of AAS are calcium

silicate hydrate (C–S–H) with a low Ca/Si ratio,

hydrotalcite and Aluminate-Ferrite-mono(sulphate) hy-

drate phase (AFm).10,11 AFm phases have the general

formula [Ca2(Al,Fe)(OH)6].X.xH2O, where X denotes

one formula unit of a single charged anion, or half a

formula unit of a doubly charged anion. The C–S–H

phases in OPC and alkali-activated slags show differ-

ences in morphology, crystallinity and chemical com-

position.10,12,13

In the OPC, topotactic growth of C–S–H is pre-

dominant, whereas in the AAS C–S–H is formed in

the pore space.10,11,13 Different silicate concentrations

in the pore solution are at least partially responsible for

these differences.14 Slag activated by Na-metasilicate

exhibits a relatively high concentration of Si in solution

and forms a foil-like C–S–H phase (as observed with

transmission electron microscopy12), whereas OPC has

a lower Si concentration and has a needle-like C–S–H

phase. C–S–H with a low Ca/Si ratio, as in AAS, is

also characterised by a higher Al/Ca ratio than C–S–H

in the OPC system.15,16

In this study the hydration process and the develop-

Advances in Cement Research, 2006, 18, No. 3, July, 119–128

119

0951-7197 # 2006 Thomas Telford Ltd

* EMPA, Materials Science and Technology, Uberlandstrasse 129,

8600 Dubendorf, Switzerland.

(ACR 5526) Paper received 31 May 2005; accepted 12 April 2006

ment of the microstructure in an alkali-activated slag

system is investigated using a wide range of analytical

techniques and compared with the well known OPC

system.

Experiment

Materials

Experiments were carried out with a slag without

activator, an alkali-activated slag (AAS) and an ordin-

ary Portland cement (CEM I 42.5 N). The AAS

consists of slag which was mixed with solid Na-

metasilicate pentahydrate (Na2SiO3.5H2O) at a ratio of

91 : 9 wt %.

The chemical composition of the slag without acti-

vator and the OPC is given in Table 1. The negative

loss on ignition of the slag was due to the oxidation of

sulphide. X-ray diffraction (XRD) analyses showed that

the slag contained approximately 10 wt % merwinite,

some minor amounts of calcite and anhydrite (due to

limited oxidation and carbonation of the slag). The

specific surface area (blaine) of the slag is 6730 cm2/g;

of the AAS 5640 cm2/g (slag + activator), and of the

OPC 2960 cm2/g. Slag had a relatively high fraction of

fine particles (, 20 �m) whereas OPC had somewhat

coarser particles.

Methods

Samples for X-ray fluorescence (XRF) were analysed

by Philips PW 2400. Mineralogical composition was

determined by XRD using a Siemens D500 powder

diffractometer. Particle size distribution was determined

with a laser granulometer, Malvern Mastersizer X.

Slag cement pastes were prepared with a water/

cement ratio (w/c) of 0.3 (OPC with a w/c of 0.5) and

mixed twice for 90 s according to EN 196-3. Calori-

metric measurements (conduction calorimeter TAM

Air) were carried out with 5 g of the fresh paste; the

remaining paste was cast in 0.5 l polyethylene bottles,

sealed and stored under controlled conditions at 208C.

Pore fluids of the hardened samples were extracted

using the steel die method17 with pressures up to

530 N/mm2, the solutions were filtered immediately

(0.45 �m nylon filter). The pH was analysed with a

combined pH-electrode which was calibrated against

KOH solutions; the total concentrations of Al, Ca, Fe,

K, Mg, Na, Si and S were determined by inductively

coupled plasma optical emission spectrometry (ICP-

OES) in samples diluted by a factor 10 with HNO3

(6.5%) to prevent the precipitation of solids. The

concentrations of SO42–, SO3

2– and HS– were also

determined by colorimetry, iodometry and ion chroma-

tography (cf. Fig. 1). In addition, Si was also measured

in more diluted (factor 100) samples (in HCl) to verify

that no Si had precipitated in the acidified samples.

A fraction of the solid paste was removed before

pore fluid extraction, crushed and submersed in acetone

to remove the pore solution, dried at 408C in an oven

and then used for thermogravimetric analysis (TGA)

and XRD analyses. TGA measurements were carried

out under N2 with powdered samples at a heating rate

of 20 K/min up to 9808C.

Samples were also examined by scanning electron

microscopy (Philips ESEM FEG XL 30) using back-

scattered, secondary electron images and energy dis-

persive X-ray (EDX) analysis of polished surfaces.

Sample preparation included pressure impregnation

with epoxy resin, cutting, polishing and then coating

with carbon.

Results and discussion

Strength development and heat evolution

Both the investigated OPC and AAS showed similar

28 day strength, however the early strength evolution of

the two systems showed significant differences. The

AAS was characterised by a higher early strength (Table

2), but after 1 day, the strength development and hydra-

tion rates decreased in the AAS system. The strength

development of slag without activator was much slower.

Calorimetric measurements (Fig. 2) show that the

heat flow evolution of the AAS system was charac-

terised by a relatively long dormant period and an

intense, but narrow peak between � 13 and � 20 h

(maximum at 16 h). This exothermal activity is compa-

tible with the observed fast strength development of

AAS during the first day. As will be shown later, it can

be correlated with the fast reaction of the small slag

particles , 2 �m (see section entitled Microstructure

below). The system ‘slag without activator’ showed a

very low heat flow with two distinct maxima at 4 and

12 h. Calorimetric data from OPC revealed a very

broad peak between 4 and 25 h with a maximum at

approximately 12 h and a shoulder at 17 h. On a

qualitative level, the broad peak and the longer exother-

mal activity reflect a slower hydration of the OPC

system leading to lower early strength. However, the

total amount of heat cannot be compared between the

three systems, since the formation of different reaction

products leads to different reaction enthalpies. In addi-

Table 1. XRF analyses of the main element oxides in wt % for pure slag without waterglass and for OPC

Oxides: wt % SiO2 Al2O3 CaO MgO Fe2O3 Na2O K2O SO3 LOI

Slag 34.8 10.7 42.5 7.7 0.92 0.09 0.38 2.4 –0.34

OPC 19.7 4.7 63.2 1.9 2.7 0.08 1.1 3.4 2.56

Gruskovnjak et al.

120 Advances in Cement Research, 2006, 18, No. 3

tion, the absolute height of the peaks should not be

compared with each other, but rather their relative

shapes should be compared.

Strength development strongly depends on the inter-

actions between solid and liquid components; that is

dissolution rates, saturation/chemical equilibrium of the

pore solution, precipitation or hydration rates and micro-

structural evolution. The following sections are thus

focused on these various aspects of cement hydration, in

order to explore the mechanisms behind the observed

strength and heat flow developments in AAS and OPC.

Hydration products

The results of TGA/differential thermal analysis

(DTG) analyses for slag and AAS at different hydration

times are shown in Fig. 3. The comparison of the TGA

spectra with measured reference spectra of the pure

cement hydrate phases (not shown in Fig. 3) indicates the

presence of C–S–H, AFm and/or hydrotalcite (Ht). The

C–S–H and Ht were also identified with XRD (see

below). In the OPC system, portlandite (CH) and ettrin-

gite were formed in addition to C–S–H and AFm.18

Because of different phase assemblages a direct compari-

son of TGA results from AAS and OPC is omitted.

The XRD analyses of AAS that are shown in Fig. 4

indicate that C–S–H was a main hydration product

with a high peak at approximately 308 2Ł (as high-

lighted in Fig. 4). This is in contrast to the OPC

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Table 2. Compressive strength development in MPa of

mortars AAS, OPC and slag without activator; with the w/c

ratios used

AAS OPC Slag

w/c ¼ 0.3: w/c ¼ 0.5: w/c ¼ 0.5:

Days MPa MPa MPa

1 21 10 n.a.

2 29 23 n.a.

7 40 38 13

28 56 50 24

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Fig. 2. Heat flow calorimetry for AAS, OPC and slag

without activator

Hydration of alkali-activated slag: comparison with OPC

Advances in Cement Research, 2006, 18, No. 3 121

system, in which usually no such distinct peak is

observed, for C–S–H, as discussed by Taylor.19 The

high peak in the AAS spectra indicates a higher

crystallinity of C–S–H in the AAS system. For AAS, a

C–S–H peak was already visible after 1 day, whereas

the pure slag reacted more slowly.

After the first day the C–S–H peak height of AAS

increased only slightly. Again, this demonstrates the

slow hydration rate after the first day. The C–S–H

formation after 1 day is either very slow or the products

are X-ray amorphous. In addition to C–S–H, peaks

from quartz and the main crystalline slag phase

(merwinite) can be identified in all samples. Merwinite

is hardly dissolved during the first 180 days. Hydro-

talcite can be identified after 28 and 180 days (128 2Ł).

Characteristics of C–S–H. EDX analyses of C–S–

H after 28 and 3–7 days of hydration for AAS and

OPC, respectively, shows a linear trend in the graph of

Ca/Si plotted against Ca/Al (Fig. 5), which indicates a

constant substitution of Si by Al.16 A similar linear

trend is observed for both AAS and OPC; however, the

slope is steeper (m ¼ 11.2) for OPC than for AAS

(m ¼ 3.4). This reflects a lower Al concentration of C–

S–H in the OPC system.

Hydrotalcite is closely mixed with C–S–H; the

uptake of Al in the C–S–H phase may be estimated

from the intercept in an Al/Ca to Mg/Ca ratio plot.20

Extrapolating the line to Mg/Ca ¼ 0 indicates that the

C–S–H contains a significant amount of Al (Al/

Ca � 0.1).

C–S–H in slag systems is characterised by elevated

crystallinity and therefore longer Si–O–Si dreierketten

chains, as documented with NMR.21 Consequently

more bridging tetrahedra are available within the C–S–

H structure. Bridging tetrahedra are the main potential

sites for Si substitution by Al.22 Thus, the higher Al

content in C–S–H from AAS in Fig. 5 may be

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Fig. 3. Thermo-gravimetric analysis (TGA) of AAS (unhydrated, 1, 7, 28 and 180 days) and slag without activator (28 days).

Upper part shows the total loss of weight in %, the lower part shows the differential of weight %. Peaks from waterglass,

C–S–H phase, hydrotalcite (Ht) and AFm phase can be observed

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Fig. 4. XRD analyses of slag without activator (unhydrated,

28 days) and AAS (unhydrated, 1, 7, 28 and 180 days) allow

a qualitative identification of the different mineral and

glassy phases. Ht, hydrotalcite; Q, quartz; M, merwinite;

C, calcite; and C–S–H phase

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Fig. 5. Ca/Si against Ca/Al atomic ratio plot of EDX

analyses of C–S–H phase in AAS and OPC (just from the

inner product IP) after a few days of hydration. The mean

Ca/Si ratios and the slopes of the two systems are indicated

Gruskovnjak et al.

122 Advances in Cement Research, 2006, 18, No. 3

correlated with a higher crystallinity of C–S–H. The

mean Ca/Si ratio of the C–S–H was very different in

the two systems: 0.8 in AAS and 1.3 in OPC. The low

Ca/Si ratios in AAS reflect the relatively high concen-

tration of Si and low Ca concentration in the pore

solution (see below). The mean Ca/(Si+Al) ratios of

0.55 for AAS and 1.18 for OPC, respectively, indicate

a higher Al substitution in AAS.

In Table 3 different atomic ratios are given from

EDX analyses of C–S–H in the three systems: AAS,

OPC and slag without activator. Whereas the AAS

shows a high Na uptake, the OPC and the pure slag

system without activator exhibit much lower Na/Ca

ratios. The Ca/Si ratio of slag without activator is

closer to that of OPC, but there is still a higher Al

incorporation observed than in the OPC system.

Pore solution chemistry

In the OPC system, the composition of the pore

solution is dominated during the first hours by K,

S(VI) and a high pH, whereas lower concentrations of

Na, Ca, Si, Al are observed.18,23 After 12–24 h,

sulphate concentrations decrease rapidly as anhydrite

and gypsum are consumed by the precipitation of

ettringite and at the same time the pH increases

accordingly. As hydration proceeds, K, Na concentra-

tions and the pH in the OPC system increase slowly as

the amount of pore solution decreases due to the

precipitation of hydrates and as more and more alkalis

(originally trapped in the clinkers) are released into the

pore solution (Table 4).18

The composition of the pore solution in the AAS

system is different (Table 4); Ca concentrations are 10

times lower than in the OPC system, whereas Si

concentrations are 50 to 100 times higher. The AAS

system is dominated by a high pH and a high Na, S

content and somewhat lower K, Al, and Si concentra-

tions. The presence of sulphide in the slag leads to

sulphide (HS–) in the pore solution. The HS– present

partially oxidises to sulphite (SO32–) and sulphate

(SO42–). Sulphate also stems from the dissolution of

anhydrite in the slag (Table 4). Sulphide oxidation is

complex as various products and intermediates are

produced. The main intermediates and products ob-

served in alkaline solutions are thiosulphate (S2O32–),

sulphite and sulphate.24–26 In the pore solutions investi-

gated, sulphide, sulphite and sulphate were determined.

Table 3. Different atomic ratios of C–S–H in AAS, OPC and

slag without activator (with EDX analysis)

Atomic ratio AAS OPC Slag

Ca/Si 0.8 1.3 1.2

Al/Ca 0.5 0.1 0.4

Na/Ca 0.5 , 0.1 , 0.1

K/Ca , 0.1 , 0.1 , 0.1

Table4.Chem

icalanalysesofthepore

solutionsextracted

from

AASlandOPC

(mmol/l)

Na:

K:

Li:

Ca:

Al:

Si:

P:

S(V

I):

S(IV):

S(–II):*

Stotaly:

pH

Sam

ple

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

mmol/l

AAS-1

day

153

102

29

1. 4

0. 25

7. 0

11

0. 06

47{

16

87

410

13. 7

AAS-7

days

133

102

38

2. 2

0. 35

4. 8

8. 5

0. 03

36{

17

350

n.a.

13. 6

AAS-28days

143

102

45

1. 2

0. 36

4. 3

9. 6

0. 03

39{

20

260

550

13. 7

AAS-180days

133

102

48

1. 2

0. 47

3. 2

5. 9

0. 03

46{

24

350

n.a.

13. 7

OPC-1

day

44

430

0. 75

4. 2

0. 17

0. 07

––

––

83}

13. 4

OPC-7

days

55

520

0. 81

2. 1

0. 24

0. 09

––

––

10}

13. 6

OPC-28days

63

560

0. 92

1. 2

0. 27

0. 12

––

––

11}

13. 6

OPC-105days

57

650

0. 86

1. 5

0. 21

0. 04

––

––

17}

13. 8

MeasuredBa,

Cr,Fe,

Mg,MoandSrconcentrationsin

theAASsystem

werebelow

therespectivedetectionlimitsBa,

Mo,Sr,

0. 01;Cr,Fe,

0. 02andMg,

0. 04mmol/l.n.a.,notanalysed.AAS:K,Ca,

Si,

AlandpH

meanoftwoindependentdeterminations.

*Relativeerrorofsulphidemeasurements

�20%;

y MeasuredbyICP-O

ESin

alkalinesolutions;

thesum

ofanalysedS(V

I),S(IV)andS(–II)is

lower

then

Stotalas

differentinorganic

andorganic

S-species

mightbepresent,whichhavenotbeenquantified,{MeasuredbyIC;

} MeasuredbyICP-O

ES

inacidifiedsolutions,

StotalcorrespondsforOPC

toS

(VI).

Hydration of alkali-activated slag: comparison with OPC

Advances in Cement Research, 2006, 18, No. 3 123

Comparison with the measured total sulphur concentra-

tions (Table 4) indicates the presence of additional

sulphur species, most probably thiosulphate. The pre-

sence of relatively high concentrations of thiosulphate

was established by Glasser et al.27 in the pore solution

of a blended slag–OPC system.

During the first hours and days, the chemical

composition of the pore solution is dominated by the

dissolution of fast-reacting components such as Na-

metasilicate and the small fraction of anhydrite and

calcite present in the slag. In addition, small slag

particles (, 2 �m) dissolve – as observed by environ-

mental scanning electron microscope (ESEM) (cf sec-

tion Microstructure) – within 7 days, which is

consistent with the presence of Al, K and sulphide in

the pore solutions (Table 4).

The dissolution of the Na-metasilicate activator leads

initially to relatively high Na and Si concentrations in

the pore solution of the AAS system (Table 4), which

decrease with time as C–S–H precipitates. Puertas et

al.28 reported for slag activated with Na-metasilicate

around 2 mol/l Na and Si in the pore solution after 3 h,

which was drastically reduced after 24 h as precipitat-

ing C–S–H also included, in addition to Ca and Si,

significant amounts of alkalis and Al (cf section

entitled Characteristic of C–S–H above, and

Refs 29–31).

In contrast to the decreasing Na concentrations, K

concentrations slowly increase with time, as K stems

from the slow dissolution of the slag particles whereas

Na originates mainly from the Na-metasilicate activa-

tor. Similarly to Na, K also fractionates between pore

solution and C–S–H. Al and Ca concentrations in the

pore solution remain more or less constant with time as

they are in equilibrium with the precipitation of solid

phases such as C–S–H, hydrotalcite or AFm.

In contrast to OPC, which is characterised by a high

Ca/Si+Al ratio in the pore solution and the presence of

portlandite, in the AAS system a low Ca/Si ratio in the

pore solution is observed, which corresponds to the

lower Ca/Si ratio in the precipitated C–S–H. The pore

solution of AAS has an approximately 10 times lower

Ca concentration than the pore solution of OPC, but a

comparable pH value. This indicates that in the AAS

system the pore solution is undersaturated with respect

to portlandite, which agrees with the absence of portlan-

dite from the solid phases in AAS. Similarly, the high

Al concentrations in the solution correspond to the

observed high Al/Si ratio in the C–S–H phase. These

findings agree with NMR studies,15,16 which have

shown that increasing amounts of Al in solution result

in a C–S–H phase with an increased quantity of Al

incorporated in the chain structure and corresponding

with an increased chain length and a higher crystallinity.

Microstructure

The microstructure of AAS is dominated by the

relatively fast formation of a gel-like C–S–H matrix in

the pore space (Fig. 6). The interstitial space between the

slag grains is already completely filled with a gel-like C–

S–H after 24 h. This interstitial matrix is somehow

comparable with the so-called outer product (OP) in OPC

(see discussion below). At higher magnifications it is

observed, that between 1 and 7 days the small slag

particles (, 2 �m) become fully hydrated. In addition,

thin layers (, 1 �m) of dense C–S–H are formed around

larger slag particles. However, these hydration layers

grow very slowly. After 7 days their thickness is still

approximately 1 to 2 �m (Fig. 7). Between 7 and 28 days

the microstructural evolution becomes very slow: the

hydration layers reach thicknesses of a few micrometres

and the interstitial gel becomes slightly denser (Fig. 8).

The microstructural evolution in AAS demonstrates fast

hydration rates during the first few days followed by

relatively slow reactions. This evolution correlates with

the relatively fast development of strength during the first

*��'���� ���

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Fig. 6. ESEM image of AAS after 1 day of hydration.

Backscattered image (high vacuum) on an impregnated,

polished and coated sample showing unhydrated slag

particles (bright) in an interstitial gel-like mass

$���������������'���� ��

/���������������

A���� �����'���� ��

Fig. 7. ESEM image of AAS after 7 days of hydration.

Backscattered image (high vacuum) on an impregnated,

polished and coated sample showing unhydrated slag

particles (bright), fully hydrated slag particles and particles

forming a hydration rim

Gruskovnjak et al.

124 Advances in Cement Research, 2006, 18, No. 3

few days and the subsequent deceleration of strength

development (Table 2).

In contrast, reaction rates in OPC do not change

drastically and strength development is more uniform

(Table 2). This also corresponds with the microstructur-

al evolution in OPC. After 1 day, a hydration layer of

approximately 2 �m has formed around the clinker

grains. The interstitial space is still predominately

empty (or rather filled with pore solution, see Fig. 9).

Only after several days is a ‘topotactic’ growth of

needle-like C–S–H observed to fill the space between

the cement grains.32 This so-called outer product (OP)

forms bridges between the single grains and is thus

responsible for the strength development. The growth

rate of the OP is kinetically controlled by the hydration

and dissolution rates of the clinker grains. The latter

controls the propagation of hydration front towards the

core of the clinker grains, that is the formation of the

inner product (IP).

As in AAS, the main hydration product in slag

without activator consists of a homogeneous interstitial

mass. As observed in ESEM, the microstructure of

hydrated slag without activator after 28 days is similar

to the microstructure of AAS after 1 day of hydration.

Furthermore, in this case the IP and OP cannot be

distinguished. The formation of the IP is so small, that

hardly any reaction rims can be identified on the

surface of the slag grains even after 28 days. These

observations fit with the measured strength develop-

ment (Table 2). The compressive strength of AAS after

1 day is similar to the strength values of slag without

activator after 28 days of hydration.

Comparison of the systems

Figures 10–12 illustrate the critical parameters

which dominate the hydration processes in AAS and in

OPC. Fig. 10 shows that the chemical composition of

the pore solutions is strongly influenced by the fast-

dissolving components of the raw materials, which are

mainly represented by Na-metasilicate in AAS and by

alkali-sulphates in OPC. The dissolution of these reac-

tive phases leads to a high alkalinity of the pore

solutions in both systems after a short time, whereby

Na is predominant in AAS and K in OPC. In AAS the

Si and Al concentrations of the pore solutions are

higher than in OPC, whereas the Ca concentration is

relatively low. In addition, AAS is undersaturated with

respect to portlandite and ettringite. The high Si

concentrations in AAS are related to the fast dissolu-

tion of Na-metasilicate, which is an important factor

for the voluminous precipitation of C–S–H within the

interstitial space already during the first 24 h. Slag

hydration is faster in systems activated with Na-

metasilicate than in systems activated with KOH,

NaOH or Ca(OH)2.33 This illustrates the importance of

Na-metasilicate for the early strength development of

AAS cements (for our AAS: 21 MPa after 1day).

In addition to the Na-metasilicate component, the

high fineness of the slag represents a further important

factor for the fast reactivity of AAS. AAS generally

has a higher fineness than the OPC in order to achieve

a comparable early strength (AAS-cement with a blaine

value of 5644 cm2/g; OPC with a blaine of 2960 cm2/

g). The AAS cement thus contains a much greater

amount of small particles. As determined by laser

granulometry, 22 vol % of the slag component in the

AAS cement is represented by particles with diameters

below 2 �m. These small particles are completely

dissolved or hydrated within the first 24 h, whereas

hydration of the larger particles is much slower. This

fits well with the results of thermodynamic modelling

(see Lothenbach and Gruskovnjak34), which indicate

that 20 vol % of the glass component is dissolved after

$���������������'���� ���

A������������'���� ���

/�������������

Fig. 8. ESEM image of AAS after 28 days of hydration.

Backscattered image (high vacuum) on an impregnated,

polished and coated sample showing unhydrated slag

particles (bright), fully hydrated slag particles and particles

forming a hydration rim. The interstitial mass is much

denser than after 1 day of hydration

���1���������

/������������

5������� �

Fig. 9. ESEM image of OPC after 1 day of hydration.

Backscattered image (high vacuum) on an impregnated,

polished and coated sample showing unhydrated clinker

particles (bright), black ¼ porosity and dark grey ¼hydration product

Hydration of alkali-activated slag: comparison with OPC

Advances in Cement Research, 2006, 18, No. 3 125

1 day. After 7 and 28 days, only 23 and 27 vol %

respectively have reacted;34 this corresponds with the

observed decrease of the strength evolution after the

first days of hydration.

The reduced strength development rate indicates that

the hydration mechanisms are changing with time. To a

lesser extent, this is also the case in the OPC system.

As soon as the reactive components (alkali-sulphates in

OPC, Na-metasilicate and small slag particles in AAS)

are consumed, the hydration processes are kinetically

controlled by the more slowly dissolving phases, which

are represented by major components of both systems

(glass and merwinite in AAS, C3S/C3A/C2S/C4AF in

OPC). In addition, the IP forms a protective rim around

the larger grains, which potentially hinders ion ex-

change between the unhydrated cores and the pore

solution. The formation of the IP thus acts as an

additional kinetic barrier.

In contrast to the slow dissolution process, the

precipitation of the hydrate phases is relatively fast

once saturation is reached. Consequently chemical

equilibrium of hydrates with pore solution is already

established within the first day. Because of the fast

precipitation, changes of the pore solution chemistry

are controlled by the equilibrium with the hydrate

phases in both systems. However, the differences of the

pore solution chemistry lead to very different hydrate

assemblages as shown in Fig. 11: the AAS system is

undersaturated with respect to portlandite and ettringite.

In addition, also the C–S–H phases exhibit very differ-

ent chemical and morphological characteristics in the

two systems. In AAS the C–S–H is characterised by

much lower Ca/Si ratios, higher incorporation of Na

and higher substitution of Al for Si. The latter can be

correlated with a higher crystallinity, which results in a

higher amount of suitable lattice sites (bridging tetra-

hedra) for Al substitution.

In addition, the microstructural evolutions in the AAS

and OPC systems exhibit several important differences,

as illustrated in Fig. 12. First, the formation of an IP can

hardly be observed in AAS during the first 28 days and

hence the concept of IP and OP cannot be applied for

AAS in the same way as for OPC. The fast formation of a

gel-like interstitial matrix in AAS is obviously the

product of a ‘through solution’ precipitation that con-

trasts with the slower mechanism of ‘topotactic’ growth

of C–S–H, which dominates the formation of OP in

OPC. The different growth mechanisms also lead to the

different morphological and structural characteristics of

the C–S–H phase: in SEM studies of AAS pastes the C–

S–H appears as a gel-like matrix whereas in OPC the C–

S–H is characterised by a needle-like morphology. The

gel-like morphology in AAS could be misinterpreted as a

less well defined hydration product than the needle-like

C–S–H in OPC. However, as deduced from the XRD

spectra, the crystallinity of the C–S–H is higher in AAS.

��� ��1������������)��%-/,

���:

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5����������� ������

-/

%*

��

(�

%-/

(�

*#�"�

��

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/������������ �

)����$'����,

""* -5�

��������� ����)������,

'�����)��� ����������,

���:��#�B���

��1��)��*#���*#����#��.$,

""* -5�����=�

"�

�� ����������������

:���'�����)��(�-/,

��������� ��

Fig. 10. Comparison of AAS and OPC hydration; summar-

ising the reactivity of solid components in the cement and

the resulting pore solution chemistry

/������������ �

""* -5�

�C*C/ �C*C/

"$�

������� �

���'�="$�

��������

����� ��� ������C*C/

""* -5�

������

D���=*0���

"��� ��������

(�0� ��������

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C

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Fig. 11. Comparison of AAS and OPC hydration; summarising the difference in the hydrate assemblages and the

characteristics of the C–S–H phase

Gruskovnjak et al.

126 Advances in Cement Research, 2006, 18, No. 3

This is compatible with a foil-like structure that is

observed for C–S–H in AAS at higher magnifications

when using transmission electron microscopy.12 It is

suggested that the ‘topotactic’ growth mechanism in

OPC only results in a nanoscopic short-range order of

several atom layers and hence the mesoscopic C–S–H

needles do not reflect euhedral shapes of single C–S–H

crystals.

Conclusions

The development of strength in AAS and OPC pastes

from the same strength category shows significant

differences. A much higher early strength and a slow

increase between 7 and 28 days are typical for the AAS

system in contrast to a more continuous strength

development in OPC. This behaviour can be attributed

to different hydration mechanisms and to components

of different reactivity.

(a) The early hydration of AAS is characterised by the

high reactivity of Na-metasilicate and the relatively

fast dissolution of the smallest fraction of slag

particles (, 2 �m) which leads to a high Si

concentration in the pore solution and to saturation

with respect to C–S–H.

(b) As a consequence, AAS is dominated by the fast

precipitation of a gel-like C–S–H phase, which fills

the former pore space already after 1 day; in contrast,

the needle-like outer product C–S–H in OPC only

forms bridges between the clinker grains slowly due

to topotactic growth. This leads to a slower increase

of the strength in OPC.

(c) The distinct compositions of the C–S–H phases

reflect the chemical characteristics of the pore

solutions in the different systems. In contrast to the

OPC system, the C–S–H phase as well as the pore

solution in AAS exhibit low Ca/Si ratios. For AAS, a

high incorporation of Al and Na in C–S–H has been

observed.

(d) After a few days the hydration mechanisms and

kinetics in AAS change as the fast-dissolving small

slag particles are consumed. The process is then

kinetically controlled by the larger and more slowly

dissolving slag particles. In addition the hydration

rims, which cover the surface of the slag grains, may

act as a kinetic barrier for diffusion processes.

However, in contrast to the OPC system, in which the

unhydrated cores of the cement grains are protected

by a thick shell of inner product, the protective layers

in the AAS system are very thin. This may explain

the more distinct increase of strength in AAS from 7

to 28 days.

Acknowledgements

The authors would like to thank Luigi Brunetti for

preparation of numerous samples, Claudia Schreiner

and Oliver Nagel for pore water analyses and valuable

hints regarding sample preservation, Urs Gfeller and

Christoph Zwicky for XRD and XRF analyses, Josef

Kaufmann for measurements of the grain size distribu-

tion and Walther Trindler and his team for the help in

the laboratory.

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& ����� �����

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3���������� =35

35���������""*����-5�

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1 January 2007

Gruskovnjak et al.

128 Advances in Cement Research, 2006, 18, No. 3