# Effect of sulfur on the deformation mechanisms of zirconium alloy at intermediate temperatures (25-400°C)

Franck Ferrer, Alain Barbu, Thierry Bretheau, Jérôme Crépin

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Franck Ferrer, Alain Barbu, Thierry Bretheau, Jérôme Crépin. Effect of sulfur on the deformation mechanisms of zirconium alloy at intermediate temperatures (25-400°C). 15th International Conference on Structural Mechanics in Reactor Technology (SMiRT-15), Aug 1999, Séoul, South Korea. pp.469-476. hal-00116554

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# Effect of Sulfur on the Deformation Mechanisms of Zirconium Alloys at Intermediate Temperatures (25-400° C)

Franck Ferrer¹), Alain Barbu²), Thierry Bretheau¹) and Jérôme Crepin¹)

1) Laboratoire de Méchanique des Solides, France

2) Laboratoire des Solides Irradiés, France

**ABSTRACT** : Cezus, France, recently showed that a small addition of sulfur in α-zirconium and its alloys could reduce steady state creep-rate by three at 400°C under a 130MPa applied stress. Sulfur surely acts directly on the deformation controlling mechanisms. But these medium temperature deformation mechanisms are still not well understood. The present study aims both at the understanding of sulfur effect and at the determination of the controlling mechanisms in this temperature range. The work is conducted on ZrNbwt1% and on α-zirconium samples containing different sulfur concentrations. The study consists in a comparative analysis of these materials at three different scales : global macroscopic scale, SEM scale and TEM scale.

## INTRODUCTION

Cezus, France, recently showed that a small addition of sulfur of about 20 wt ppm reduces steady state creep rate by 3 at 400°C under a 130MPa applied stress (figure 1) (1). A so strong effect of a so low sulfur concentration suggests that sulfur directly acts on the deformation mechanisms. Nevertheless these mechanisms are still not well understood although a lot of studies have been performed (2-5). The present study aims both at the understanding of the effect of sulfur and at a better knowledge of the controlling mechanisms in this temperature range, specially under creep conditions.

As very little is known about the intermediate temperature deformation mechanisms, all potential elementary mechanisms are considered : grain boundary sliding (GBS), self-diffusion and dislocation glide (cross-slip, lattice friction). Each one is then discussed with regards to experimental results and to the plausible effects of a very low amount of sulfur in order to get the relevant one(s). This will in return enable a better understanding of the effect of sulfur.

## Materials and Experimental Techniques :

Internal pressure creep tests are conducted on two ZrNb1wt% alloys (Table 1) containing different sulfur amount ([S]<5ppm and [S]=25ppm). We make sure that sulfur does not induce microstructural changes : in both materials there is the same rolling texture and the mean grain size is about 4µm ; the precipitates of βNb have the same chemical composition, the same size and the same density.

Constant strain-rate tensile tests are performed on α-zirconium plates ([S]<5ppm and [S]=30ppm) containing 1200 wt ppm of oxygen. The texture is also a rolling texture and the mean grain size is about 50µm (material A). In order to test the effect of grain size, some samples are submitted to thermo-mechanical treatments ; we thus obtain material B with a
medium mean grain size of 200µm and material C with a large mean grain size of 800µm. In these two latter cases the texture remains unchanged.

<figure><img src="image_2.png" /><figcaption>Figure 1: Creep curves of two ZrNb1wt % containing different sulfur concentrations.</figcaption></figure>

### Domain of Existence of the Sulfur Effect:

The sulfur effect has been pointed out in creep tests at 400°C under a 130MPa applied stress on all industrial zirconium alloys (1). It is first interesting to observe the sulfur effect dependence on temperature and stress under creep conditions. The creep tests are performed on pre-pressurized tubular specimens which are placed in an electric furnace. Creep strain is defined as the variation $\Delta D/D$ of the tube outer diameter. These creep tests are performed on ZrNb1wt% tubes for temperatures between 350°C and 430°C and hoop stresses between 90 and 150MPa. The effect is observed in all the temperature and stress ranges with a creep rate ratio of about 3 between the high and low sulfur content materials. The secondary creep rate follows in both materials a classical law of the type $\dot{\epsilon}_s=A\sigma^n\exp(-Q/RT)$ with $Q \approx 215$ kJ/mol and $n \approx 4$ (figure 2).

Tensile tests under constant strain rate ($\dot{\epsilon}=6.10^{-5} s^{-1}$) are also performed on α-zirconium (material A) with high and low sulfur content between room temperature and 500°C. The sulfur effect appears at $T=100°C$, increases and reaches a maximum for $T=400°C$ (figure 3). Let us now consider the different potential controlling mechanisms.

<table><caption>Table 1: Chemical composition of the ZrNb1wt % alloy.</caption><thead><tr><th>Element</th><th>Nb (wt %)</th><th>O (wt ppm)</th><th>Fe (wt ppm)</th><th>H (wt ppm)</th><th>Others (wt ppm)</th></tr></thead><tbody><tr><td>Analysis</td><td>1.01</td><td>1200</td><td>108</td><td>&lt;5</td><td>&lt;50</td></tr></tbody></table>

## CONTROLLING MECHANISM

### Grain Boundary Sliding (GBS):

Grain boundary sliding may act in creep either as a deformation mechanism or as a controlling mechanism. In the first case, creep deformation is mainly given by GBS; in the second case, deformation is given by another deformation process (dislocation glide for instance) and GBS is needed to accommodate deformation incompatibilities between grains.

To test the case of GBS as a deformation mechanism fiducial micro grids (6) are deposited prior to deformation on the sample surface (ZrNb1 % alloy). The samples are then submitted to creep at 430°C under 130MPa up to 5% hoop strain. Figure 4 shows that, in both materials with and without sulfur, there is no shearing of the grid lines when crossed by a grain
boundary. It implies that there is no major grain boundary sliding inside the plane of the
observed surface. Nevertheless grid deformation shows a 5% deformation in the hoop
direction. Thus, GBS is not the mechanism giving the main deformation ; nevertheless it may
control creep rate by enabling the accommodation of grain boundaries incompatibilities. In such
a case very little GBS may be needed and the previous technique is inadequate for the detection.

<figure><img src="image_3.png" /><figcaption>Figure 2 : $\dot{\epsilon}_s$ vs T in an Arrhenius plot (a) and $\dot{\epsilon}_s$ vs $\sigma$ in a ln/ln plot (b).</figcaption></figure>

To check whether GBS is the controlling mechanism or not, we could observe how the
effect of sulfur depends on the mean grain size. We then perform tensile tests under constant
strain rate ($\dot{\epsilon}$=6.10$^{-5}$s$^{-1}$) at 400°C on materials A,B and C (figure 5 (a)). These results show that
the strengthening amount $\Delta\sigma(\varepsilon)$ due to sulfur is independent of the grain size. Figure 5 (b)
shows that both materials, with high and low sulfur content, follow a Hall Petch law of the
type $\sigma=\sigma_0+kd^{-1/2}$; it appears clearly that sulfur only modifies the $\sigma_0$ value, which is precisely
independent of grain size, and not the k value which describes the grain boundary effect. This
leads to two conclusions : 1. sulfur is not acting in the grain boundaries and 2. GBS is not the
controlling mechanism in this temperature range (25°C-400°C). We now consider lattice self
diffusion as a potential controlling mechanism.

<figure><img src="image_4.png" /><figcaption>Figure 3 : $\sigma_{2\%}$ vs T for the two $\alpha$-zirconium plates (d=50µm) containing different sulfur concentrations.</figcaption></figure>

Lattice Self Diffusion :

Lattice self-diffusion is often evoked as being the creep controlling mechanism at intermediate
temperatures (7) since it controls dislocation climb. In such a case, the apparent creep activation
energy Q generaly corresponds to the self-diffusion enthalpy $\Delta H_{sd}$. $\Delta H_{sd}$ is not well known in zirconium and appears to strongly depend on the impurity content (e.g Fe). Thus, the Q value adequation is not a determining argument.

<figure><img src="image_5.png" /><figcaption>Figure 4: Grids after 400h creep at T=430°C (a) [S]=25 ppm and (b) [S]< 5 ppm.</figcaption></figure>

To go further, we may wonder whether self-diffusion is high enough, in this temperature range, to justify the measured creep rates. Let us consider a simple model (8) of bulk self-diffusion controlled creep which leads to the following expression for the creep rate :

$$ \dot{\epsilon}_s^c = A \frac{\mu b}{kT} D_L \left( \frac{\sigma}{\mu} \right)^n \quad (1) $$

where $A \approx 10$, $D_L$ is the lattice self diffusion coefficient, $\mu$ the shear modulus and $n=4$.

<figure><img src="image_6.png" /><figcaption>Figure 5: Grain size effect at T=400°C and $\dot{\epsilon}=6.6 \cdot 10^{-5} s^{-1}$. (a) Stress-strain curves for materials A and C and (b) $\sigma_{1\%}$ (T=400°C, $\dot{\epsilon}=6.6 \cdot 10^{-5} s^{-1}$) as a function of the grain size d.</figcaption></figure>

Taking $D_L=10^{-27} m^2 s^{-1}$ (the maximum value found in literature (9)) at T=400°C and $\sigma=130$ Mpa, Eq.(1) gives $\dot{\epsilon}_s^c=8,5 \cdot 10^{-13} s^{-1}$. This value is 4 orders of magnitude lower than the experimental creep rates $\dot{\epsilon}_s=1,5 \cdot 10^{-8} s^{-1}$. Whatever the details of a model based on lattice self-diffusion, it will underestimate the experimental creep rate. This implies that lattice self-diffusion cannot be the controlling mechanism in this temperature range.

Moreover it can be mentioned that the sulfur effects appear at 100°C which is a much too low temperature for a significant bulk diffusion.

Then bulk diffusion is not a relevant controlling mechanism at intermediate temperatures and
the sulfur effect does not result from a sulfur/vacancy interaction.

We have shown that neither grain boundary sliding nor lattice self diffusion are the mechanisms controlling the deformation kinetics. Sulfur is then necessarily acting on dislocations and the controlling mechanism in this temperature range is necessarily a mechanism involving dislocations (cross-slip, lattice friction, jogs dragging, core self-diffusion...). Before going further we need to have a better understanding of the sulfur-dislocation interaction.

## SULFUR-DISLOCATION INTERACTION

Before we focus more precisely on the sulfur-dislocation interaction, we first may wonder if sulfur does modify or not the nature of the slip systems.

### Slip Systems Activated in Both Materials with High and Low Sulfur Content:

We identify the activated slip systems in materials A thanks to two complementary techniques : analysis of slip lines traces at the sample surface and analysis of dislocations in a Transmission Electron Microscope. In both techniques the identification is performed after 2% tensile strain at the constant strain rate $\dot{\epsilon}=6.10^{-5}s^{-1}$.

*At T≤200°C :*

The slip line analysis reveals that, in both materials, prismatic slip is the slip system mainly activated (90%) ; other slip systems are exceptional (1st order pyramidal slip). Besides, the slip lines are very straight and well defined (figure 6 (a)). Thin foils cut parallel to the basal plane show long straight screw <a> dislocations. In general two systems are activated in a single grain (figure 7 (a)). The presence of these screw dislocations means that in this temperature range the screw dislocations are submitted to a strong Peierls force as in titanium (10). This Peierls force is surely due to the dislocation core structure which is non planar in zirconium and titanium (11). No difference is observed between the two materials with high or low sulfur content.

*At T>200°C :*

The results are quite different even if concerning the sulfur effect we still observe no difference between the high and low sulfur materials. Prismatic slip remains the mainly activated slip system (80%) but 1st order pyramidal slip{1101}<1120> is now more frequently observed (15%). In addition the slip lines become curved or wavy, often near the grain boundaries, indicating that cross-slip has occurred (figure 6 (b)). In some cases it is possible to identify the cross-slip plane (figure 6 (b)). The main slip trace then always corresponds to a prismatic one and the cross-slip trace to a pyramidal plane. Then cross-slip from prismatic to 1st order pyramidal planes occurs for T>200°C in high and low sulfur materials.

Cross-slip is also put in evidence by TEM observations. Only <a> dislocations are present but this time they are bowed in pyramidal planes (figure 7 (b)). Long prismatic loops and dipoles can be observed in thin foils cut parallel to prismatic planes ; their identification shows that they surely result from a double cross-slip mechanism as in Ti (10). Both slip line analysis and TEM observations show that cross-slip from prismatic planes to 1st order pyramidal planes is occurring in both materials for temperatures higher than 200°C under constant strain rate tensile tests.

Dislocations generated under creep conditions in the ZrNb1% alloy show the same kind of microstructure constituted by an homogeneous distribution of <a> type dislocations bended in pyramidal planes ; they undoubtedly are the signature of cross-slip. Even if the slip traces are not observable on the sample surface because of the surface orientation and oxidation it is highly probable that prismatic slip remains the principal activated slip system under creep
conditions at 400°C. It results that cross-slip from prismatic to pyramidal planes is also occurring during creep.

It then appears that sulfur does not modify the activated slip systems in the temperature range 25°C-400°C. Prismatic slip is the main slip system in the whole temperature range and 1st order pyramidal slip and cross-slip to 1st order pyramidal planes appear at temperatures higher than 200°C.

<figure><img src="image_7.png" /><figcaption>Figure 6: Slip lines after 2% tensile strain at 25°C (a) and at 400°C (b). The A direction corresponds to a prismatic plane and the B to one of the two associated pyramidal planes.</figcaption></figure>

### Dynamic Strain Ageing or... :

Concerning the sulfur-dislocation interaction, three different cases are possible : 1. sulfur is immobile, 2. sulfur is mobile and follows the dislocations (Dynamic Strain Ageing (DSA)) or 3. sulfur atoms segregate into the dislocation core and modify it. Since the <a> screw dislocation core is non planar (10) we may imagine that if sulfur segregates it will modify the core structure : all the parameters attached to dislocation are then affected (Peierls force, cross-slip, core self-diffusion...).

### Sulfur is immobile :

Friedel (12) showed that the hardening due to solid solution atoms homogeneously distributed in the matrix is given by :

$$ \Delta\sigma = \frac{c.E_i}{2b^3} \qquad (2) $$

where c is the impurity concentration, Eᵢ the dislocation-impurity interaction energy and b the burgers vector of the dislocations. Taking Eᵢ=1eV (an upper bound for this kind of interaction energy), c=45ppm and b=3,2Å in Eq.(2) leads to Δσ=0.1MPa. This value is considerably too low compared to the observed hardening (25MPa at 400°C) ; this is due to the too low sulfur concentration c. Moreover if the sulfur was immobile the effect would be effective at low temperature and would decrease with increasing temperature because of thermal activation. This is not what is observed ; sulfur is then not immobile and we must rather consider cases 2 and 3.

### DSA or segregation :

The purpose of this discussion is to know whether sulfur segregates in the dislocation cores or if sulfur interacts as a classical solute atmosphere and we do not aim at describing the nature of the interaction between sulfur atoms and dislocations (elastic, electronic...).

In case 2 (DSA) the effect of sulfur is only present in a well defined temperature and strain rate range. At a given temperature T sulfur atoms diffuse at a velocity vₛ. If dislocations are
moving with a constant velocity $v_d>>v_s$, sulfur cannot follow the dislocations and there is no dragging force. If dislocations are moving too low, $v_d<<v_s$, the equilibrium is always instantaneously reached around the dislocations and no dragging force is exerted. The dragging force then only exists when $v_d\approx v_s$ where a dynamic equilibrium is reached.

<figure><img src="image_8.png" /><figcaption>Figure 7: Dislocation microstructure after 2% tensile strain at 25°C (a) and at 400°C (b). In both cases the plane of observation corresponds to the basal plane.</figcaption></figure>

In case 3 the situation is quite different for $v_d<<v_s$. Under these conditions sulfur atoms still move very fast but this time it means that, as they can always follow dislocations, they can segregate and always stay into dislocation cores: the core is then continuously modified and the effect will still be present even if $v_d<<v_s$.

Figure 3 shows that on material A the effect reaches a maximum at $T=400°C$ at a strain rate of $\dot{\epsilon}=6.10^{-5} s^{-1}$; under these conditions we then surely have $v_d\approx v_s$. We perform tests at the same temperature but at a much lower constant strain rate $\dot{\epsilon}=2.10^{-8} s^{-1}$ on the two materials A; under these conditions we surely have $v_d<<v_s$. Figure 8 (a) shows that the effect is still present and is as strong as at a higher strain rate ($\dot{\epsilon}=6.10^{-5} s^{-1}$); segregation is then much more probable than DSA effects. Besides, we perform tests at 250°C at a high strain rate ($\dot{\epsilon}=5.10^{-2} s^{-1}$) so that this time $v_d>>v_s$. Figure 8(b) shows that the effect of sulfur tends to disappear: now the velocity of dislocations is too high and the atoms of sulfur cannot follow and stay in the dislocation core.

## BACK TO THE CONTROLLING MECHANISM...

We have shown that sulfur atoms probably segregate into the dislocation cores and modify their geometry. Concerning the controlling mechanism the question remains open. Indeed once the core of the dislocations is modified all the parameters concerning the dislocations are modified: cross-slip which implies core recombination, dislocation mobility, core self-diffusion... However Pujol (13) has shown that there is a critical stress $\sigma_c$ in creep at 200°C: $\sigma<\sigma_c$ leads to exhausted creep and $\sigma>\sigma_c$ leads to secondary creep and then to rupture. A better understanding of this phenomenon and the possible effect of sulfur on $\sigma_c$ will surely enable us to have a better knowledge of the controlling mechanism. Such tests are in progress.

## CONCLUSION

This study shows that neither lattice self-diffusion nor GBS can be the creep controlling
mechanism for temperatures lower than 430°C. An effect of sulfur through an interaction with vacancy mobility or grain boundaries is as well eliminated. Sulfur then surely interacts with dislocations which implies that the controlling mechanism involves dislocations (cross-slip, lattice friction, core diffusion...). The sulfur concentration is very low which excludes an interaction as an immobile impurity. Besides we have shown that the effect of sulfur is present in a very large range of temperatures and strain rates : this surely excludes an effect of dynamic strain ageing. It is rather suggested that sulfur segregates into the dislocation cores and modifies its geometry. This assumption is now to be confronted to numerical simulations of dislocation cores in the presence of sulfur.

<figure><img src="image_9.png" /><figcaption>Figure 8 : Stress-strain curves for materials A at $\dot{\epsilon}=2.10^{-8}s^{-1}$(a) and at $\dot{\epsilon}=6.10^{-1}s^{-1}$ (b).</figcaption></figure>

Concerning the controlling mechanism we should say that it is a mechanism which involves dislocations : cross-slip, lattice friction, core diffusion. Complementary tests are needed to get a better knowledge of the controlling mechanism.

Acknowledgments : The authors are grateful to Framatome and Cezus for their financial support and for permitting to publish this work. They also wish to thank Zircotube for conducting the creep tests.

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