The damping measurements were carried out in a vacuum (about 30 Pa) at room temperature. The specimens fixed at one end were excited into resonance by a permanent magnet fixed at the free side of the bending beam and a sinusoidal alternating magnetic field. The internal friction was characterised by the logarithmic decrement δ of the free decay of the vibrating beam δ = ln (An/An+1), where An and An+1 are the amplitudes of a free decay of vibrations after n and (n+1) cycles, respectively. The resonance frequency ranged from 130 to 160 Hz. The strain amplitude dependences of the logarithmic decrement were measured. The specimens were annealed step by step at increasing temperatures up to 500 °C for 0.5 h and after each annealing cycle quenched into water of room temperature. The annealing at higher temperatures was performed in an argon atmosphere to avoid oxidation. The damping measurements were carried out immediately after heat treatment and quenching at room temperature.

Figure 23 shows the logarithmic decrement plotted against the logarithm of the maximum strain amplitude for μ-Mg+3nZrO2 samples. The logarithmic decrement was measured before (as received - as rec) and after step by step annealing at increasing upper temperature of the cycle. It can be seen that the measured strain dependence of the logarithmic decrement exhibits two regions and can be expressed as

0 H δ ε = δ +) δ( (ε), (8)

where δ0 is the amplitude independent component (or only weakly dependent on the maxi‐ mum strain amplitude) found in the first region, for lower strain amplitudes. In the second region, for higher strain amplitudes, the component δH(ε) depends on the strain amplitude (amplitude dependent internal friction – ADIF); it increases with increasing strain amplitude. The upper temperature of the thermal cycling influences the logarithmic decrement. In the case of the μ-Mg+3nZrO2 composite, the thermal treatment influences mainly the amplitude independent component of the decrement. It is important to note that very high values of -Mg+3nZrO2

amplitude

can be expressed as

0 H δ ε( ( ) = δ + δ ε), (8)

temperature.

0.030 as rec

100°C 200°C 300°C 400°C 500°C

decrement  0.000 0.005 0.010 0.015 0.020 0.025

500 °C (b).

depends on the kind of particles.

damping in the amplitude independent component - in order of 10-2 - were obtained. The strain amplitude dependences of the logarithmic decrement for μ-Mg+3nAl2O3 are given in Figures 24. Figure 24a shows the strain amplitude dependence of the logarithmic decrement for specimens annealed up to 300 ºC. Figure 24b shows the strain amplitude dependence of the logarithmic decrement for specimens annealed at temperatures between 300 and 500 ºC. The thermal cycling influences mainly the amplitude dependent component δH that increases with increasing upper temperature of the cycle, if the upper temperature is between room temper‐ ature and 300 ºC. For the upper temperature between 300 and 500 ºC, the δH component decreases with increasing upper temperature. where δ0 is the amplitude independent component (or only weakly dependent on the maximum strain amplitude) found in the first region, for lower strain amplitudes. In the second region, for higher strain amplitudes, the component δH(ε) depends on the strain amplitude (amplitude dependent internal friction – ADIF); it increases with increasing strain amplitude. The upper temperature of the thermal cycling influences the logarithmic decrement. In the case of the μ-Mg+3nZrO2 composite, the thermal treatment influences mainly the amplitude independent component of the decrement. It is important to note that very high values of damping in the amplitude independent component - in order of 10-2 - were obtained. The strain amplitude dependences of the logarithmic decrement for μ-Mg+3nAl2O3 are given in Figures 24. Figure 24a shows the strain amplitude dependence of the logarithmic decrement for specimens annealed up to 300 ºC. Figure 24b shows the strain amplitude dependence of the logarithmic decrement for specimens annealed at temperatures between 300 and 500 ºC. The thermal cycling influences mainly the amplitude

The strain amplitude dependences of the logarithmic decrement measured in UFG-Mg after thermal cycling are given in Figure 25a. Only a slight influence of thermal cycling on the δ versus ε curves can be observed. Figure 25b shows the plots of δ versus ε for UFG-Mg+3nGr. It can be seen that the thermal cycling causes a slight decrease of the logarithmic decrement at small strain amplitudes. The δH component changes only very slightly with increasing temperature of the cycle. Figure 26 shows the strain amplitude dependence of the logarithmic decrement for various materials before annealing (as received). The value of the critical strain εc at which the logarithmic decrement begins to increase with strain amplitude depends on

Mechanical and Acoustic Properties of Magnesium Alloys Based (Nano) Composites Prepared by Powder…

decrement δ

a) b)

**Figure 25.** Strain amplitude dependence of the decrement for UFG-Mg (a) and UFG-Mg+3nGr (b), annealed at in‐

It is generally accepted that the *δ*0 component is connected with the material microstructure and the amplitude dependent component *δ*H is due to an interaction of dislocations with point defects (e.g. solute atoms). The dislocations contribute to the damping capacity by the motion of vibrating dislocation lines. Particulates (Al2O3, ZrO2, Gr), as secondary phases, may improve the damping capacity of alloys. The particles in a composite may cause an increase in the dislocation density as a result of thermal strain mismatch between the ceramic particles and the matrix during preparation and/or thermal treatment. The difference between the coeffi‐ cients of thermal expansion (CTE) of the particles and the matrix may create the thermal residual stresses after cooling from the processing temperature to room temperature. The thermal stresses are also generated during heat treatment and when composites are heated and cooled through a temperature range (thermal cycling). In generally, there are biaxial or triaxial stress fields located near the reinforcement-matrix interface. The amplitudes of the stress field decrease with increasing distance from the interface. In a simple one-dimensional model, the thermal stresses *σ*TS, produced by a temperature change ΔT, at the interface are

0.010

0.015

0.020

0.025

0.030

0.035

UFG-Mg+3nGr

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187

as rec 100°C 200°C 300°C 400°C 450 °C

amplitude ε 10-5 10-4 10-3

3

annealing temperature and it depends on the kind of particles.

UFG-Mg

amplitude ε

10-5 10-4 10-3

as rec. 100 °C 200 °C 300 °C 400 °C 500 °C

decrement δ

0.00

Figure 25

creasing temperature.

given by (Chawla, 1991)

0.01

0.02

0.03

dependent component δH that increases with increasing upper temperature of the cycle, if the upper temperature is between room

into water of room temperature. The annealing at higher temperatures was performed in an argon atmosphere to avoid oxidation.

Figure 23 shows the logarithmic decrement plotted against the logarithm of the maximum strain amplitude for -Mg+3nZrO2 samples. The logarithmic decrement was measured before (as received - as rec) and after step by step annealing at increasing upper temperature of the cycle. It can be seen that the measured strain dependence of the logarithmic decrement exhibits two regions and

The damping measurements were carried out immediately after heat treatment and quenching at room temperature.

**Figure 23.** Amplitude dependence of decrement measured after annealing at increasing temperatures.

Figure 23.Amplitude dependence of decrement measured after annealing at increasing temperatures.

Figure 24.Amplitude dependence of the decrement for -Mg+3nAl2O3 annealed at temperatures up to 300°C (a) and temperatures between 300 and **Figure 24.** Amplitude dependence of the decrement for μ-Mg+3nAl2O3 annealed at temperatures up to 300°C (a) and temperatures between 300 and 500 °C (b).

The strain amplitude dependences of the logarithmic decrement measured in UFG-Mg after thermal cycling are given in Figure 25a. Only a slight influence of thermal cycling on the δ versus ε curves can be observed. Figure 25b shows the plots of δ versus ε for UFG-Mg+3nGr. It can be seen that the thermal cycling causes a slight decrease of the logarithmic decrement at small strain amplitudes. The δH component changes only very slightly with increasing temperature of the cycle. Figure 26 shows the strain amplitude dependence of the logarithmic decrement for various materials before annealing (as received). The value of the critical strain εc at which the logarithmic decrement begins to increase with strain amplitude depends on annealing temperature and it The strain amplitude dependences of the logarithmic decrement measured in UFG-Mg after thermal cycling are given in Figure 25a. Only a slight influence of thermal cycling on the δ versus ε curves can be observed. Figure 25b shows the plots of δ versus ε for UFG-Mg+3nGr. It can be seen that the thermal cycling causes a slight decrease of the logarithmic decrement at small strain amplitudes. The δH component changes only very slightly with increasing temperature of the cycle. Figure 26 shows the strain amplitude dependence of the logarithmic decrement for various materials before annealing (as received). The value of the critical strain εc at which the logarithmic decrement begins to increase with strain amplitude depends on annealing temperature and it depends on the kind of particles.

damping in the amplitude independent component - in order of 10-2 - were obtained. The strain amplitude dependences of the logarithmic decrement for μ-Mg+3nAl2O3 are given in Figures 24. Figure 24a shows the strain amplitude dependence of the logarithmic decrement for specimens annealed up to 300 ºC. Figure 24b shows the strain amplitude dependence of the logarithmic decrement for specimens annealed at temperatures between 300 and 500 ºC. The thermal cycling influences mainly the amplitude dependent component δH that increases with increasing upper temperature of the cycle, if the upper temperature is between room temper‐ ature and 300 ºC. For the upper temperature between 300 and 500 ºC, the δH component

where δ0 is the amplitude independent component (or only weakly dependent on the maximum strain amplitude) found in the first region, for lower strain amplitudes. In the second region, for higher strain amplitudes, the component δH(ε) depends on the strain amplitude (amplitude dependent internal friction – ADIF); it increases with increasing strain amplitude. The upper temperature of the thermal cycling influences the logarithmic decrement. In the case of the μ-Mg+3nZrO2 composite, the thermal treatment influences mainly the amplitude independent component of the decrement. It is important to note that very high values of damping in the amplitude independent component - in order of 10-2 - were obtained. The strain amplitude dependences of the logarithmic decrement for μ-Mg+3nAl2O3 are given in Figures 24. Figure 24a shows the strain amplitude dependence of the logarithmic decrement for specimens annealed up to 300 ºC. Figure 24b shows the strain amplitude dependence of the logarithmic decrement for specimens annealed at temperatures between 300 and 500 ºC. The thermal cycling influences mainly the amplitude dependent component δH that increases with increasing upper temperature of the cycle, if the upper temperature is between room temperature and 300 ºC. For the upper temperature between 300 and 500 ºC, the δH component decreases with increasing upper

into water of room temperature. The annealing at higher temperatures was performed in an argon atmosphere to avoid oxidation.

Figure 23 shows the logarithmic decrement plotted against the logarithm of the maximum strain amplitude for -Mg+3nZrO2 samples. The logarithmic decrement was measured before (as received - as rec) and after step by step annealing at increasing upper temperature of the cycle. It can be seen that the measured strain dependence of the logarithmic decrement exhibits two regions and

The damping measurements were carried out immediately after heat treatment and quenching at room temperature.

amplitude 10-5 10-4 10-3

> decrement

a) b) Figure 24.Amplitude dependence of the decrement for -Mg+3nAl2O3 annealed at temperatures up to 300°C (a) and temperatures between 300 and

The strain amplitude dependences of the logarithmic decrement measured in UFG-Mg after thermal cycling are given in Figure 25a. Only a slight influence of thermal cycling on the δ versus ε curves can be observed. Figure 25b shows the plots of δ versus ε for UFG-Mg+3nGr. It can be seen that the thermal cycling causes a slight decrease of the logarithmic decrement at small strain amplitudes. The δH component changes only very slightly with increasing temperature of the cycle. Figure 26 shows the strain amplitude dependence of the logarithmic decrement for various materials before annealing (as received). The value of the critical strain εc at which the logarithmic decrement begins to increase with strain amplitude depends on annealing temperature and it

**Figure 24.** Amplitude dependence of the decrement for μ-Mg+3nAl2O3 annealed at temperatures up to 300°C (a) and

0.000 0.002 0.004 0.006 0.008 0.010 0.012

300°C 400°C 500°C

amplitude 10-5 10-4 10-3

decreases with increasing upper temperature.

can be expressed as

0 H δ ε( ( ) = δ + δ ε), (8)

temperature.

0.030 as rec

100°C 200°C 300°C 400°C 500°C

decrement  0.000 0.005 0.010 0.015 0.020 0.025

500 °C (b).

decrement

186 Light Metal Alloys Applications

amplitude 10-5 10-4 10-3

> decrement

depends on the kind of particles.

0.000 0.002 0.004 0.006 0.008 0.010 0.012 0.000

Figure 23.Amplitude dependence of decrement measured after annealing at increasing temperatures.

as rec. 100°C 150°C 200°C 250°C 300°C

temperatures between 300 and 500 °C (b).

amplitude 10-5 10-4 10-3

0.005

0.010

0.015

0.020

0.025

0.030 as rec

100°C 200°C 300°C 400°C 500°C

**Figure 23.** Amplitude dependence of decrement measured after annealing at increasing temperatures.

Figure 25 **Figure 25.** Strain amplitude dependence of the decrement for UFG-Mg (a) and UFG-Mg+3nGr (b), annealed at in‐ creasing temperature.

It is generally accepted that the *δ*0 component is connected with the material microstructure and the amplitude dependent component *δ*H is due to an interaction of dislocations with point defects (e.g. solute atoms). The dislocations contribute to the damping capacity by the motion of vibrating dislocation lines. Particulates (Al2O3, ZrO2, Gr), as secondary phases, may improve the damping capacity of alloys. The particles in a composite may cause an increase in the dislocation density as a result of thermal strain mismatch between the ceramic particles and the matrix during preparation and/or thermal treatment. The difference between the coeffi‐ cients of thermal expansion (CTE) of the particles and the matrix may create the thermal residual stresses after cooling from the processing temperature to room temperature. The thermal stresses are also generated during heat treatment and when composites are heated and cooled through a temperature range (thermal cycling). In generally, there are biaxial or triaxial stress fields located near the reinforcement-matrix interface. The amplitudes of the stress field decrease with increasing distance from the interface. In a simple one-dimensional model, the thermal stresses *σ*TS, produced by a temperature change ΔT, at the interface are given by (Chawla, 1991)

3

$$
\sigma\_{\rm TS} = \frac{\rm E\_{\rm F} E\_{\rm M}}{\left(E\_{\rm F} \rm f + E\_{\rm M} \left(1 - \rm f\right)\right)} \rm f \Delta a \Delta T \tag{9}
$$

difference in CTEs for both components of the composite. In mc and UFG Mg no new dislo‐ cations are created during thermal cycling. Hence, it is expected only a slight influence of the thermal cycling on the δ versus ε curves, which is in agreement with experimental results.

Mechanical and Acoustic Properties of Magnesium Alloys Based (Nano) Composites Prepared by Powder…

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189

decrement

0.00

Plastic zones may also contribute to the damping behaviour.

(Zhang et al., 1994a)

0.01

0.02

0.03

0.04

amplitude 10-5 10-4 10-3

**Figure 26.** Amplitude dependence of the decrement for μ-Mg, UFG-Mg and μ-Mg and UFG-Mg reinforced with np.

Interfaces may contribute to the damping behaviour of composite materials. Due to a large difference in the coefficients of thermal expansion between the particles and the matrix, high thermal residual stresses are created around the particles. The thermal stresses with temper‐ ature can relax and plastic zones of tangled dislocation are formed. Their average radius depends on the difference in CTEs, particles radius and the change of temperature. The higher the matrix yield stress, the smaller radius of the plastic zones (Dunand & Mortensen, 1991).

The contribution of interfaces to internal friction depends on the character of interfacial bonding. The interface damping should be considered. The internal friction of a composite with the perfectly bonded interface (interface is incoherent) depends on the shape of particles, their volume fraction and the magnitude of local stress at the particle-matrix interface. Assuming that all the particles have the same diameter, the logarithmic decrement is given as

> 4.5(1 ) (2 ) n

 n

where *V*p is the volume fraction of particles and *ν* is the Poisson ratio of the matrix. The effect of the perfectly bonded interfaces on the damping behaviour becomes more significant at high temperatures; in Mg based composites this could be above 150 ºC. The matrix becomes

p

where *Δα* is the mismatch of the CTE values between the reinforcements and the matrix, *E*F and *E*M are the values of the Young's modulus for the reinforcement and the matrix, respec‐ tively and f is the volume fraction of the reinforcement. The CTE of the matrix is much higher than that of the reinforcement. Thus after cooling, the tensile thermal residual stresses are created in the matrix. The thermal stresses may relax around the matrix-reinforcement interface by emitting dislocation. An increase in the dislocation density near reinforcement has been calculated as (Arsenault & Shi, 1986, Dunand & Mortensen, 1991):

$$
\Delta \varphi = \frac{B \int \Delta a \, \Delta T}{b \, (\mathbf{l} - \mathbf{f}')} \frac{\mathbf{l}}{\mathbf{t}} \tag{10}
$$

where t is the minimum dimension of reinforcement, b is the magnitude of the Burgers vector of dislocations and *B* is a geometrical constant (depending on the aspect ratio). The newly formed dislocations may be sources of the higher damping capacity because of the motion of vibrating dislocation lines under cycling loading. According to the Granato-Lücke theory (Granato, & Lücke, 1981), weak and strong pinning points restrict motion of dislocations in their glide plane. At low strain amplitudes (stresses), dislocation bows out between the weak pinning points, leading to the amplitude independent logarithmic decrement. The dislocation remains anchored at the strong pinning points. At higher strain amplitudes, the force exerted on the weak pinning points becomes higher than the binding force of the weak pinning points. The dislocation segments break away from some weak pinning points within longer disloca‐ tion segments. This leads to a drastic instantaneous increase in the dislocation strain and thus giving rise to the level of the logarithmic decrement. The amplitude dependent component of the decrement depends on the dislocation density, length of shorter and longer dislocation segments. At temperatures higher than 0 K, the breakaway of shorter dislocation segments is thermally activated. During thermal cycling, new dislocations are created due to the difference in the CTEs (see Eq. (10)). An increase in the dislocation density, while the number of weak pinning points remains constant, increases effectively the length of dislocation segment between weak pinning points and hence the δH component should increase, which is observed. The higher temperature of the cycling, the higher the dislocation density and the higher increase in the effective length of the shorter dislocation segments. The critical strain εc at which the decrement becomes to be amplitude dependent is determined by the interaction energy between dislocations and the weak pinning points and the distance between the pinning points. As the distance between weak pinning point increases with the upper temperature, the dislocations can break away from the weak pinning point under lower critical stress and thus, the value of εc should decrease with increasing upper temperature, which is observed. The values of εc for a composite are proportional to the yield stress of the composite matrix. It should be consider that the value of εc can be influenced by the thermal stresses due to the difference in CTEs for both components of the composite. In mc and UFG Mg no new dislo‐ cations are created during thermal cycling. Hence, it is expected only a slight influence of the thermal cycling on the δ versus ε curves, which is in agreement with experimental results.

( ( )) F M

Ef E 1 f

= DaD

E E f T

where *Δα* is the mismatch of the CTE values between the reinforcements and the matrix, *E*F and *E*M are the values of the Young's modulus for the reinforcement and the matrix, respec‐ tively and f is the volume fraction of the reinforcement. The CTE of the matrix is much higher than that of the reinforcement. Thus after cooling, the tensile thermal residual stresses are created in the matrix. The thermal stresses may relax around the matrix-reinforcement interface by emitting dislocation. An increase in the dislocation density near reinforcement has

1

(1 ) a

where t is the minimum dimension of reinforcement, b is the magnitude of the Burgers vector of dislocations and *B* is a geometrical constant (depending on the aspect ratio). The newly formed dislocations may be sources of the higher damping capacity because of the motion of vibrating dislocation lines under cycling loading. According to the Granato-Lücke theory (Granato, & Lücke, 1981), weak and strong pinning points restrict motion of dislocations in their glide plane. At low strain amplitudes (stresses), dislocation bows out between the weak pinning points, leading to the amplitude independent logarithmic decrement. The dislocation remains anchored at the strong pinning points. At higher strain amplitudes, the force exerted on the weak pinning points becomes higher than the binding force of the weak pinning points. The dislocation segments break away from some weak pinning points within longer disloca‐ tion segments. This leads to a drastic instantaneous increase in the dislocation strain and thus giving rise to the level of the logarithmic decrement. The amplitude dependent component of the decrement depends on the dislocation density, length of shorter and longer dislocation segments. At temperatures higher than 0 K, the breakaway of shorter dislocation segments is thermally activated. During thermal cycling, new dislocations are created due to the difference in the CTEs (see Eq. (10)). An increase in the dislocation density, while the number of weak pinning points remains constant, increases effectively the length of dislocation segment between weak pinning points and hence the δH component should increase, which is observed. The higher temperature of the cycling, the higher the dislocation density and the higher increase in the effective length of the shorter dislocation segments. The critical strain εc at which the decrement becomes to be amplitude dependent is determined by the interaction energy between dislocations and the weak pinning points and the distance between the pinning points. As the distance between weak pinning point increases with the upper temperature, the dislocations can break away from the weak pinning point under lower critical stress and thus, the value of εc should decrease with increasing upper temperature, which is observed. The values of εc for a composite are proportional to the yield stress of the composite matrix. It should be consider that the value of εc can be influenced by the thermal stresses due to the

D D D = - *Bf T*

+ - (9)

*b ft* (10)

F M

been calculated as (Arsenault & Shi, 1986, Dunand & Mortensen, 1991):

r

TS

s

188 Light Metal Alloys Applications

**Figure 26.** Amplitude dependence of the decrement for μ-Mg, UFG-Mg and μ-Mg and UFG-Mg reinforced with np.

Interfaces may contribute to the damping behaviour of composite materials. Due to a large difference in the coefficients of thermal expansion between the particles and the matrix, high thermal residual stresses are created around the particles. The thermal stresses with temper‐ ature can relax and plastic zones of tangled dislocation are formed. Their average radius depends on the difference in CTEs, particles radius and the change of temperature. The higher the matrix yield stress, the smaller radius of the plastic zones (Dunand & Mortensen, 1991). Plastic zones may also contribute to the damping behaviour.

The contribution of interfaces to internal friction depends on the character of interfacial bonding. The interface damping should be considered. The internal friction of a composite with the perfectly bonded interface (interface is incoherent) depends on the shape of particles, their volume fraction and the magnitude of local stress at the particle-matrix interface. Assuming that all the particles have the same diameter, the logarithmic decrement is given as (Zhang et al., 1994a)

$$\delta = \frac{4.5(1-\nu)}{\pi(2-\nu)} V\_p \tag{11}$$

where *V*p is the volume fraction of particles and *ν* is the Poisson ratio of the matrix. The effect of the perfectly bonded interfaces on the damping behaviour becomes more significant at high temperatures; in Mg based composites this could be above 150 ºC. The matrix becomes relatively soft relative to the nanoparticles. Hence, the internal friction at the interface may be thermally activated.

In the case of weak bonding at the interface, interfacial slip (sliding at the interface) may occur. In this case, the frictional energy loss caused by the sliding at interfaces may become a primary source of damping. The damping component due to interfacial slip under the applied stress amplitude *σ*0 can be expressed as (Zhang et al., 1994a)

$$\delta = \frac{3\pi^2}{2} \frac{\kappa \sigma\_r (\varepsilon\_0 - \varepsilon\_{cr})}{\sigma\_0^2 / E\_C} V\_p \tag{12}$$

**4.2. Temperature relaxation spectrum of internal friction**

tan

f/Hz

0.01

**Figure 28.** Arrhenius plots estimated for both peaks.

0.1

1

10

0.00

0.02

0.04

0.06

0.08

0.10

amplitude was 1.2x10-4. The heating as well as cooling rate was 1 K/min.

0.5 Hz

**Figure 27.** Temperature relaxation spectrum of IF estimated for various frequencies.

0.1 Hz

The temperature dependent internal friction (TDIF), internal friction spectrum, is the temper‐ ature or frequency dependence of damping. Internal friction spectra were measured in a DMA 2980 (TA Instruments) apparatus in a single heating and cooling process. Measurements were performed at various frequencies from 0.1 up to 2 Hz. Throughout the measurements the strain

Mechanical and Acoustic Properties of Magnesium Alloys Based (Nano) Composites Prepared by Powder…

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temperature (°C) 0 100 200 300 400 500 600

1/TP (K-1) 0.0010 0.0015 0.0020 0.0025 0.0030

DH=1.77 eV DH=1.16 eV

2 Hz

1 Hz

where κ is the friction coefficient between both components of the composite, *σ*r is the radial stress at the particle-matrix interface corresponding to the stress amplitude *σ*0, *ε*0 is the strain amplitude corresponding to *σ*0. The critical interface strain εcr corresponding to the critical interface shear stress τcr is the strain at which the sliding on the on interface begins. *E*C is the elastic modulus of the composite. For weakly bonded interfaces, the critical strain *ε*cr is assumed to be much lower than the strain amplitude *ε*0 and then

$$
\delta \mathcal{S} = \frac{3\pi^2}{2} \frac{\kappa \sigma\_r}{\sigma\_0} V\_p. \tag{13}
$$

The stress concentration factor *σ*r/*σ*0 has been reported to be 1.1-1.3 (Zhang et al., 1994b). The model does not take into account possible effects of temperature or frequency on damping and therefore the predictions of the model may be taken only as the first approximation. Very high values of δ0 for mc and nc Mg reinforced by zirconia np and Gr np are very probably caused by interfacial slip due to weak bonding between particles and matrix. The effect of an addi‐ tional damping due to the influence of particle can be clearly seen in Figure 26. Therefore, the effect of an additional damping due to the influence of particles may be estimated. The measured increase in *δ*0 due to the addition of ZrO2 np to μ-Mg and Gr np to UFG-Mg is about 0.008-0.009 (Figure 26). Taking for the friction coefficient a typical value of ~0.08, the relation (10) predicts a value for an additional damping owing to interfaces of *δ*0i = 0.04. The discrepancy between the predicted and measured value may be caused by a non-uniform strain state in specimens because they were subjected to bending in experiment. Hence, the stress (strain) level can reach its critical value at which the interface sliding starts only in some region of the measured specimen. Accordingly, the measured interface-damping component is lower than the predicted value. High values of the decrement component *δ*0 obtained for mc and UFG-Mg are very probably caused by grain boundary sliding support by diffusion processes. On the other hand, the observed decrease in the amplitude independent component of μ-Mg +3nAl2O3 comparing with μ-Mg may be explained according to the following way: Np situated in grain boundaries prevent the grain boundary sliding which contribute significantly to the amplitude independent component of decrement.

#### **4.2. Temperature relaxation spectrum of internal friction**

relatively soft relative to the nanoparticles. Hence, the internal friction at the interface may be

In the case of weak bonding at the interface, interfacial slip (sliding at the interface) may occur. In this case, the frictional energy loss caused by the sliding at interfaces may become a primary source of damping. The damping component due to interfacial slip under the applied stress

> 0 2 0 3 ( ) 2 /

*C*

where κ is the friction coefficient between both components of the composite, *σ*r is the radial stress at the particle-matrix interface corresponding to the stress amplitude *σ*0, *ε*0 is the strain amplitude corresponding to *σ*0. The critical interface strain εcr corresponding to the critical interface shear stress τcr is the strain at which the sliding on the on interface begins. *E*C is the elastic modulus of the composite. For weakly bonded interfaces, the critical strain *ε*cr is

*V*

*E* (12)

= *r Vp* (13)

s

2

0 3 . 2 p ks

The stress concentration factor *σ*r/*σ*0 has been reported to be 1.1-1.3 (Zhang et al., 1994b). The model does not take into account possible effects of temperature or frequency on damping and therefore the predictions of the model may be taken only as the first approximation. Very high values of δ0 for mc and nc Mg reinforced by zirconia np and Gr np are very probably caused by interfacial slip due to weak bonding between particles and matrix. The effect of an addi‐ tional damping due to the influence of particle can be clearly seen in Figure 26. Therefore, the effect of an additional damping due to the influence of particles may be estimated. The measured increase in *δ*0 due to the addition of ZrO2 np to μ-Mg and Gr np to UFG-Mg is about 0.008-0.009 (Figure 26). Taking for the friction coefficient a typical value of ~0.08, the relation (10) predicts a value for an additional damping owing to interfaces of *δ*0i = 0.04. The discrepancy between the predicted and measured value may be caused by a non-uniform strain state in specimens because they were subjected to bending in experiment. Hence, the stress (strain) level can reach its critical value at which the interface sliding starts only in some region of the measured specimen. Accordingly, the measured interface-damping component is lower than the predicted value. High values of the decrement component *δ*0 obtained for mc and UFG-Mg are very probably caused by grain boundary sliding support by diffusion processes. On the other hand, the observed decrease in the amplitude independent component of μ-Mg +3nAl2O3 comparing with μ-Mg may be explained according to the following way: Np situated in grain boundaries prevent the grain boundary sliding which contribute significantly to the

s

2

p ks

assumed to be much lower than the strain amplitude *ε*0 and then

amplitude independent component of decrement.

thermally activated.

190 Light Metal Alloys Applications

amplitude *σ*0 can be expressed as (Zhang et al., 1994a)

The temperature dependent internal friction (TDIF), internal friction spectrum, is the temper‐ ature or frequency dependence of damping. Internal friction spectra were measured in a DMA 2980 (TA Instruments) apparatus in a single heating and cooling process. Measurements were performed at various frequencies from 0.1 up to 2 Hz. Throughout the measurements the strain amplitude was 1.2x10-4. The heating as well as cooling rate was 1 K/min.

**Figure 27.** Temperature relaxation spectrum of IF estimated for various frequencies.

**Figure 28.** Arrhenius plots estimated for both peaks.

The internal friction in a material is a result of internal processes, which occurs during alternating stress cycles imposed on it and these processes originate from the interactions among structural components in the material. The presence of point, line, and planar defects within a stressed material often causes the internal friction to occur because of the atomic movement, rearrangement, or realignment of the defects under the application of a stress. These techniques can be used to characterize types of structure defects.

where k is the Boltzmann constant, T is the absolute temperature, A, B, and Cb are constants. After subtracting the background by using a fitting program PeakFit, the peak temperature TP was estimated for the both relaxation peaks. The temperatures of both peaks estimated for various frequencies are given in Table 7. The peak widths are broader than that for a Debye peak, characterised

Mechanical and Acoustic Properties of Magnesium Alloys Based (Nano) Composites Prepared by Powder…

tan

a) b)

Figure 28 shows the semilogarithmic plot of the frequency versus reciprocal value of peaks temperatures TP (estimated for both peaks) so called Arrhenius plots. From the slope and intercept of the straight line, the mean activation enthalpy for the low

Figure 28 shows the semilogarithmic plot of the frequency versus reciprocal value of peaks temperatures TP (estimated for both peaks) so called Arrhenius plots. From the slope and intercept of the straight line, the mean activation enthalpy for the low temperature peak ΔH

As far as the low temperature peak may have a dislocation nature, prestraining of the sample at low temperature should influence

As far as the low temperature peak may have a dislocation nature, prestraining of the sample at low temperature should influence the height of the peak. Therefore, the samples were de‐ formedbyrolling.Deformationbyrollingatroomtemperature (*ε*≈2.55%)increases the strength of the low temperature peak, as it is obvious from Figure 29a (the peak was extracted from the internal friction spectrum subtracting the background damping). However, the peak is stron‐ ger only in the first run after prestraining. In the second and third run, the height of the peak is approximately the same as in the as-prepared state, and the height is lower than in the first run. The strain amplitude dependent component of the logarithmic decrement (Figure 25a) indicates free dislocation loops in the material. IF peaks estimated in the second and third heating and cooling runs are practically the same. Beside this original dislocation population, the prestraining of the sample produced new dislocations which were practically completely recovered during the heating course of measurements. Owing to small grain size the grain boundaries were very probably the sources as well as sinks of the newly created dislocations. As mentioned above, the dominant deformation mode in magnesium is the basal slip; glide of  dislocations. The secondary conservative slip may be realised by motion of  dislocations on prismatic and pyramidal planes of the first kind. Couret & Caillard (1984a,b) studied by TEM prismatic glide in magnesium over a wide temperature range. They have reported that

increases the strength of the low temperature peak, as it is obvious from Figure 29a (the peak was extracted from the internal friction spectrum subtracting the background damping). However, the peak is stronger only in the first run after prestraining. In the second and third run, the height of the peak is approximately the same as in the as-prepared state, and the height is lower than

The strain amplitude dependent component of the logarithmic decrement (Figure 25a) indicates free dislocation loops in the material. IF peaks estimated in the second and third heating and cooling runs are practically the same. Beside this original dislocation population, the prestraining of the sample produced new dislocations which were practically completely recovered during the heating course of measurements. Owing to small grain size the grain boundaries were very probably the sources as well as sinks of the newly created dislocations. As mentioned above, the dominant deformation mode in magnesium is the basal slip; glide of a dislocations. The secondary conservative slip may be realised by motion of a dislocations on prismatic and pyramidal planes of the first kind. Couret & Caillard (1984a,b) studied by TEM prismatic glide in magnesium over a wide temperature range. They have reported that the screw dislocations with the Burgers vector of 1/ 3 1120 are able to glide on prismatic planes and their mobility is much lower than the mobility of edge dislocations. The deformation is controlled by the thermally activated glide of those screw dislocation segments. A single controlling mechanism has been identified as the Friedel-Escaig mechanism. This mechanism assumes a dissociated dislocation on a compact plane (0001) that joints together along a critical length *L*r producing double kinks on non-compact plane. Dislocations in very small grains are pinned at the grain boundaries and this is the reason why this dislocation population is so stable. On the other hand, the newly created dislocations during cold rolling

The high temperature peak was observed at a temperature of ≈ 345 °C (0.5Hz). The cold prestraining of the sample did not affect the strength of the high temperature peak, as it can be seen from Figure 29b. Position and height of the peak are very stable during heating as well as cooling. The activation energy was obtained from the frequency dependence of the peak temperature (Arrhenius

and their mobility is much lower than the mobility of edge dislocations. The deformation is controlled by the thermally activated glide of those screw dislocation segments. A single controlling mechanism has been identified as the Friedel-Escaig mechanism. This mechanism assumes a dissociated dislocation on a compact plane (0001) that joints together along a critical length *L*r producing double kinks on non-compact plane. Dislocations in very small grains are

The presence of the high temperature peak in magnesium has been reported by Kê (Kê, 1947). He described this peak as being

the height of the peak. Therefore, the samples were deformed by rolling. Deformation by rolling at room temperature (

**Figure 29.** Background subtracted low (a) and high (b) temperature peaks after a prestraining of 2.55%.

0.000

0.004

0.008

0.012

0.016

is the mean relaxation time,

1000/T (K-1) 1.2 1.4 1.6 1.8 2.0 2.2

¯0 are able to glide on prismatic planes

http://dx.doi.org/10.5772/57454

1 run 2 run

0 is the pre-exponential factor,

193

2.55%)

by a single relaxation time. The internal friction peak appears at the condition = 1 (Nowick & Berry. 1972) , with

1 run

2 run

3 run

is the angular frequency (= 2f, f is the measuring frequency),

1000/T [K-1] 1.8 2.0 2.2 2.4 2.6 2.8 3.0

Figure 29.Background subtracted low (a) and high (b) temperature peaks after a prestraining of 2.55%.

has been obtained and its value is 1.16±0.05 eV.

are only slightly pinned and they are recoverable during heating part of the measurement cycle.

the screw dislocations with the Burgers vector of 1 / 3 112

temperature peak H has been obtained and its value is 1.160.05 eV.

*4.2.1. Low temperature peak*

0 *wt wt DH T* exp ( ) /k , (15)

and *H* is the mean activation enthalpy.

tan

**4.2.1. Low temperature peak**

**4.2.2. High temperature peak**

plot) to be 1.770.05 eV (see Figure 28).

related to the grain boundary relaxation.

in the first run.

0.000

0.004

0.008

0.012

0.016

where

Figure 27 shows the temperature dependence of the internal friction IF=tanφ (Nowick & Berry. 1972) measured at 0.1 Hz during heating (heating rate 1 K/min.). Two peaks showing different intensities have been found: the weak low temperature peak (P1) and the more intense high temperature peak (P2). It can be seen that both peaks shift to a higher temperatures with increasing frequency, indicating that they may be relaxations peaks.

**Table 7.** Temperatures of the relaxation peaks estimated for various frequencies.

The curves introduced in Fig. 27 were measured during one heating course in the multifrequent mode. The internal friction peaks are assumed to be imposed by a background IFb expressed by

$$\text{IF}\_b = \text{A} + \text{B} \exp\left(-\text{C}\_b/\text{kT}\right),\tag{14}$$

where k is the Boltzmann constant, T is the absolute temperature, A, B, and Cb are constants. After subtracting the background by using a fitting program PeakFit, the peak temperature TP was estimated for the both relaxation peaks. The temperatures of both peaks estimated for various frequencies are given in Table 7. The peak widths are broader than that for a Debye peak, characterised by a single relaxation time. The internal friction peak appears at the condition ωτ = 1 (Nowick & Berry. 1972), with

$$\text{tot} = \text{ot}\_0 \exp\left(DH/kT\right),\tag{15}$$

where *ω* is the angular frequency (= 2πf, f is the measuring frequency), *τ* is the mean relaxation time, *τ*0 is the pre-exponential factor, and Δ*H* is the mean activation enthalpy.

where k is the Boltzmann constant, T is the absolute temperature, A, B, and Cb are constants. After subtracting the background by using a fitting program PeakFit, the peak temperature TP was estimated for the both relaxation peaks. The temperatures of both peaks estimated for various frequencies are given in Table 7. The peak widths are broader than that for a Debye peak, characterised

by a single relaxation time. The internal friction peak appears at the condition = 1 (Nowick & Berry. 1972) , with

Figure 29.Background subtracted low (a) and high (b) temperature peaks after a prestraining of 2.55%. **Figure 29.** Background subtracted low (a) and high (b) temperature peaks after a prestraining of 2.55%.

#### **4.2.1. Low temperature peak**  *4.2.1. Low temperature peak*

0 *wt wt DH T* exp ( ) /k , (15)

where

in the first run.

**4.2.2. High temperature peak**

plot) to be 1.770.05 eV (see Figure 28).

related to the grain boundary relaxation.

The internal friction in a material is a result of internal processes, which occurs during alternating stress cycles imposed on it and these processes originate from the interactions among structural components in the material. The presence of point, line, and planar defects within a stressed material often causes the internal friction to occur because of the atomic movement, rearrangement, or realignment of the defects under the application of a stress.

Figure 27 shows the temperature dependence of the internal friction IF=tanφ (Nowick & Berry. 1972) measured at 0.1 Hz during heating (heating rate 1 K/min.). Two peaks showing different intensities have been found: the weak low temperature peak (P1) and the more intense high temperature peak (P2). It can be seen that both peaks shift to a higher temperatures with

**f (Hz) P1 (°C) P2 (°C)**

0.1 79.6 313.2 0.5 94.7 344.5 1 103.3 357.9 2 109.5 367.2

The curves introduced in Fig. 27 were measured during one heating course in the multifrequent mode. The internal friction peaks are assumed to be imposed by a background IFb

where k is the Boltzmann constant, T is the absolute temperature, A, B, and Cb are constants. After subtracting the background by using a fitting program PeakFit, the peak temperature TP was estimated for the both relaxation peaks. The temperatures of both peaks estimated for various frequencies are given in Table 7. The peak widths are broader than that for a Debye peak, characterised by a single relaxation time. The internal friction peak appears at the

where *ω* is the angular frequency (= 2πf, f is the measuring frequency), *τ* is the mean relaxation

0

time, *τ*0 is the pre-exponential factor, and Δ*H* is the mean activation enthalpy.

w w

( ) b b IF = A + B exp -C /kT , (14)

*t t DH T* = exp ( / k ,) (15)

These techniques can be used to characterize types of structure defects.

increasing frequency, indicating that they may be relaxations peaks.

**Table 7.** Temperatures of the relaxation peaks estimated for various frequencies.

condition ωτ = 1 (Nowick & Berry. 1972), with

expressed by

192 Light Metal Alloys Applications

Figure 28 shows the semilogarithmic plot of the frequency versus reciprocal value of peaks temperatures TP (estimated for both peaks) so called Arrhenius plots. From the slope and intercept of the straight line, the mean activation enthalpy for the low temperature peak H has been obtained and its value is 1.160.05 eV. Figure 28 shows the semilogarithmic plot of the frequency versus reciprocal value of peaks temperatures TP (estimated for both peaks) so called Arrhenius plots. From the slope and intercept of the straight line, the mean activation enthalpy for the low temperature peak ΔH has been obtained and its value is 1.16±0.05 eV.

As far as the low temperature peak may have a dislocation nature, prestraining of the sample at low temperature should influence the height of the peak. Therefore, the samples were deformed by rolling. Deformation by rolling at room temperature (increases the strength of the low temperature peak, as it is obvious from Figure 29a (the peak was extracted from the internal friction spectrum subtracting the background damping). However, the peak is stronger only in the first run after prestraining. In the second and third run, the height of the peak is approximately the same as in the as-prepared state, and the height is lower than As far as the low temperature peak may have a dislocation nature, prestraining of the sample at low temperature should influence the height of the peak. Therefore, the samples were de‐ formedbyrolling.Deformationbyrollingatroomtemperature (*ε*≈2.55%)increases the strength of the low temperature peak, as it is obvious from Figure 29a (the peak was extracted from the internal friction spectrum subtracting the background damping). However, the peak is stron‐ ger only in the first run after prestraining. In the second and third run, the height of the peak is approximately the same as in the as-prepared state, and the height is lower than in the first run.

2.55%)

The strain amplitude dependent component of the logarithmic decrement (Figure 25a) indicates free dislocation loops in the material. IF peaks estimated in the second and third heating and cooling runs are practically the same. Beside this original dislocation population, the prestraining of the sample produced new dislocations which were practically completely recovered during the heating course of measurements. Owing to small grain size the grain boundaries were very probably the sources as well as sinks of the newly created dislocations. As mentioned above, the dominant deformation mode in magnesium is the basal slip; glide of a dislocations. The secondary conservative slip may be realised by motion of a dislocations on prismatic and pyramidal planes of the first kind. Couret & Caillard (1984a,b) studied by TEM prismatic glide in magnesium over a wide temperature range. They have reported that the screw dislocations with the Burgers vector of 1/ 3 1120 are able to glide on prismatic planes and their mobility is much lower than the mobility of edge dislocations. The deformation is controlled by the thermally activated glide of those screw dislocation segments. A single controlling mechanism has been identified as the Friedel-Escaig mechanism. This mechanism assumes a dissociated dislocation on a compact plane (0001) that joints together along a critical length *L*r producing double kinks on non-compact plane. Dislocations in very small grains are pinned at the grain boundaries and this is the reason why this dislocation population is so stable. On the other hand, the newly created dislocations during cold rolling are only slightly pinned and they are recoverable during heating part of the measurement cycle. The strain amplitude dependent component of the logarithmic decrement (Figure 25a) indicates free dislocation loops in the material. IF peaks estimated in the second and third heating and cooling runs are practically the same. Beside this original dislocation population, the prestraining of the sample produced new dislocations which were practically completely recovered during the heating course of measurements. Owing to small grain size the grain boundaries were very probably the sources as well as sinks of the newly created dislocations. As mentioned above, the dominant deformation mode in magnesium is the basal slip; glide of  dislocations. The secondary conservative slip may be realised by motion of  dislocations on prismatic and pyramidal planes of the first kind. Couret & Caillard (1984a,b) studied by TEM prismatic glide in magnesium over a wide temperature range. They have reported that the screw dislocations with the Burgers vector of 1 / 3 112 ¯0 are able to glide on prismatic planes and their mobility is much lower than the mobility of edge dislocations. The deformation is controlled by the thermally activated glide of those screw dislocation segments. A single controlling mechanism has been identified as the Friedel-Escaig mechanism. This mechanism assumes a dissociated dislocation on a compact plane (0001) that joints together along a critical length *L*r producing double kinks on non-compact plane. Dislocations in very small grains are

The high temperature peak was observed at a temperature of ≈ 345 °C (0.5Hz). The cold prestraining of the sample did not affect the strength of the high temperature peak, as it can be seen from Figure 29b. Position and height of the peak are very stable during heating as well as cooling. The activation energy was obtained from the frequency dependence of the peak temperature (Arrhenius

The presence of the high temperature peak in magnesium has been reported by Kê (Kê, 1947). He described this peak as being

pinned at the grain boundaries and this is the reason why this dislocation population is so stable. On the other hand, the newly created dislocations during cold rolling are only slightly pinned and they are recoverable during heating part of the measurement cycle.

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#### *4.2.2. High temperature peak*

The high temperature peak was observed at a temperature of ≈ 345 °C (0.5Hz). The cold prestraining of the sample did not affect the strength of the high temperature peak, as it can be seen from Fig. 29b. Position and height of the peak are very stable during heating as well as cooling. The activation energy was obtained from the frequency dependence of the peak temperature (Arrhenius plot) to be 1.77±0.05 eV (see Figure 28).

The presence of the high temperature peak in magnesium has been reported by Kê (Kê, 1947). He described this peak as being related to the grain boundary relaxation.

Grain boundaries in materials with the small grain size contain a dense network of overlapping grain boundary dislocations. Grain boundary sliding is realized by the slip and climb (pro‐ viding a maintenance of vacancy sources and sinks) of the grain boundary dislocations. Since both modes of the grain boundary dislocation motion must occur simultaneously, the slower one will control the grain boundary sliding. The climb mode involves jog formation and grain boundary diffusion and both modes of the grain boundary dislocation movement may be affected by the interaction of the grain boundary dislocations with impurities segregated in grain boundaries. According to Lakki et al. (1998) the activation energy for such process can be expressed as *ΔH* = *ΔH*j + *ΔH*GB, where *ΔH*j and ΔHGB are the activation enthalpies for the jog formation and the grain boundary diffusion, respectively. The activation enthalpy for grain boundary diffusion in the coarse grained magnesium is 0.95 eV (Frost & Ashby, 1982). Comparing with the estimated activation enthalpy value of 1.77 eV, the possible jog formation energy should be approximately 0.8 eV, which seems to be a reasonable value. Taking into consideration that the surface of grains was passivated during the milling by 1% of oxygen in a protective argon atmosphere; therefore, MgO particles in the grain boundaries may influence the grain boundary diffusion characteristics.

#### **Acknowledgements**

Z.T. and P.L. are grateful for the financial support to the Czech Grant Agency under the contract P108/12/J018.

#### **Author details**

Zuzanka Trojanová1 , Pavel Lukáč1 , Zoltán Száraz2 and Zdeněk Drozd1

1 Charles University in Prague, Czech Republic

2 The University of Manchester, UK

#### **References**

pinned at the grain boundaries and this is the reason why this dislocation population is so stable. On the other hand, the newly created dislocations during cold rolling are only slightly

The high temperature peak was observed at a temperature of ≈ 345 °C (0.5Hz). The cold prestraining of the sample did not affect the strength of the high temperature peak, as it can be seen from Fig. 29b. Position and height of the peak are very stable during heating as well as cooling. The activation energy was obtained from the frequency dependence of the peak

The presence of the high temperature peak in magnesium has been reported by Kê (Kê, 1947).

Grain boundaries in materials with the small grain size contain a dense network of overlapping grain boundary dislocations. Grain boundary sliding is realized by the slip and climb (pro‐ viding a maintenance of vacancy sources and sinks) of the grain boundary dislocations. Since both modes of the grain boundary dislocation motion must occur simultaneously, the slower one will control the grain boundary sliding. The climb mode involves jog formation and grain boundary diffusion and both modes of the grain boundary dislocation movement may be affected by the interaction of the grain boundary dislocations with impurities segregated in grain boundaries. According to Lakki et al. (1998) the activation energy for such process can

formation and the grain boundary diffusion, respectively. The activation enthalpy for grain boundary diffusion in the coarse grained magnesium is 0.95 eV (Frost & Ashby, 1982). Comparing with the estimated activation enthalpy value of 1.77 eV, the possible jog formation energy should be approximately 0.8 eV, which seems to be a reasonable value. Taking into consideration that the surface of grains was passivated during the milling by 1% of oxygen in a protective argon atmosphere; therefore, MgO particles in the grain boundaries may influence

Z.T. and P.L. are grateful for the financial support to the Czech Grant Agency under the

and Zdeněk Drozd1

, Zoltán Száraz2

and ΔHGB are the activation enthalpies for the jog

pinned and they are recoverable during heating part of the measurement cycle.

temperature (Arrhenius plot) to be 1.77±0.05 eV (see Figure 28).

He described this peak as being related to the grain boundary relaxation.

+ *ΔH*GB, where *ΔH*j

*4.2.2. High temperature peak*

194 Light Metal Alloys Applications

be expressed as *ΔH* = *ΔH*j

**Acknowledgements**

contract P108/12/J018.

**Author details**

Zuzanka Trojanová1

the grain boundary diffusion characteristics.

, Pavel Lukáč1

1 Charles University in Prague, Czech Republic

2 The University of Manchester, UK

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**Chapter 9**

), which is about 35% less than

) and more than four times lower than the density of steel

© 2014 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

**Heuristic Analysis Application for Magnesium Alloys**

The modern development of technology makes it necessary to look for new design solutions to improve efficiency and product quality, to minimize size and to increase the reliability and dimensional stability under service conditions. The choice of material is always preceded by an analysis of a number of factors including the requirements like: mechanical, physical and chemical, design, environmental, cost related, availability and weight. The problem of a too high mass verifies strongly the applicability of particular groups of materials. Therefore, over the past few decades, occurs a significant increase in demand for materials with low density and relatively high strength alloys such as titanium, aluminium and magnesium. In this group of materials special attention is devoted to magnesium alloys, which in addition to low densities have also other advantages, such as good damping capacity, the best among all currently known structural materials, high dimensional stability, good fluidity, low shrinkage of casts, a desirable combination of low density and high strength, and it can be used for machine parts operating at temperatures up to 300°C as well as relatively ease recyclability. The high damping capacity and relatively low weight allows to use magnesium alloys for production of fast moving parts where sudden changes in speed occurs. However, magnesium

**Properties Improvement**

Tomasz Tański, Krzysztof Labisz and

Additional information is available at the end of the chapter

is an undisputed advantage especially its low density (1.7 g/cm3

the density of aluminium (2.7 g/cm3

(7.86 g/cm3).

Anna Dobrzańska-Danikiewicz

http://dx.doi.org/10.5772/57587

**1. Introduction**

## **Heuristic Analysis Application for Magnesium Alloys Properties Improvement**

Tomasz Tański, Krzysztof Labisz and Anna Dobrzańska-Danikiewicz

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/57587

#### **1. Introduction**

The modern development of technology makes it necessary to look for new design solutions to improve efficiency and product quality, to minimize size and to increase the reliability and dimensional stability under service conditions. The choice of material is always preceded by an analysis of a number of factors including the requirements like: mechanical, physical and chemical, design, environmental, cost related, availability and weight. The problem of a too high mass verifies strongly the applicability of particular groups of materials. Therefore, over the past few decades, occurs a significant increase in demand for materials with low density and relatively high strength alloys such as titanium, aluminium and magnesium. In this group of materials special attention is devoted to magnesium alloys, which in addition to low densities have also other advantages, such as good damping capacity, the best among all currently known structural materials, high dimensional stability, good fluidity, low shrinkage of casts, a desirable combination of low density and high strength, and it can be used for machine parts operating at temperatures up to 300°C as well as relatively ease recyclability. The high damping capacity and relatively low weight allows to use magnesium alloys for production of fast moving parts where sudden changes in speed occurs. However, magnesium is an undisputed advantage especially its low density (1.7 g/cm3 ), which is about 35% less than the density of aluminium (2.7 g/cm3 ) and more than four times lower than the density of steel (7.86 g/cm3).

© 2014 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

#### **2. Investigation method**

Investigations described in this chapter concerns first of all investigation of structure and properties of the described engineering materials, in the state after heat and surface treatment, including the selected group of magnesium alloys, aluminium alloys and steel, as well as heuristic analysis using the methodology of integrated computer prediction, enabling the demarcation of predicted development trends of the analysed groups of materials and technologies of surface treatment and lead to the determination of their strategic development in comparison with other groups of materials and surface engineering technologies. In the initial planning stage of research scope in this investigations, in the context of choosing specific research topics with a comprehensive range of academic disciplines or specialties an innova‐ tive analysis was used to select the technology that uses the complex dendro-chronological evaluation matrix. Furthermore, in order to confirm the appropriateness of the proper selection of the investigated substrate material like the described in this paper Mg-Al-Zn cast magne‐ sium alloys, there was performed an analysis using the dendrological selection matrix, considering only three groups of most commonly used construction materials, namely steel, aluminium and magnesium alloys.

**3. Investigation results**

wide-stretching oak. (Fig. 1).

produced element.

**3.1. First part – Heuristic analysis**

In the first step of the heuristic material investigation there should be distinguished among all the available constructional materials and technologies several homogeneous groups of materials for the reason to undertake the planned experimental and comparative analysis. The impact of external positive factors (opportunities) and negative factors (difficulties) on the specific technologies analysed was evaluated using the metrological matrix of environment influence. In consistency with the concept proposed, each of the technologies assessed by the experts for their position against the macro-and microenvironment is placed in one of the quarters of this matrix. A universal scale of relative states, being a single-pole positive scale without zero, where 1 is a minimum value and 10 an extraordinarily high value, was used in the research undertaken. The dendrological matrix of values allows in a relatively easy way to visualize the results of the assessment of individual groups of materials and in a next steps of processing technologies in terms of potential, which is the real objective value of the studied area, the so-called hard values, tangible qualities as well as attractiveness, which reflects the subjective perception of its potential users,the so-called soft properties. The potential of a given group presented on the horizontal axis is the result of multi-criteria analysis carried out on the basis of expertise and extensive knowledge of literature data, taking into account the appro‐ priate proportions of the potential: creativeness, application, quality, and technical develop‐ ment. The vertical axis shows the importance of the attractiveness of the group, which is average mean value of expert and literature data evaluation in the field of research carried out on the basis of specific criteria corresponding to the economic attractiveness, cultural, eco‐ nomic, scientific and systemic. The technologies analysed were first evaluated by key experts for their attractiveness and potential and the result obtained was entered into one of the quarters of dendrological matrix of technology value. The wide-stretching oak is the most promising quarter guaranteeing the future success encompassing the technologies with a high potential and attractiveness. The soaring cypress includes very attractive technologies with a limited potential and rooted dwarf mountain pine symbolises technologies with a high potential and limited attractiveness likely to ensure a strong position for the technology if an adequate strategy is employed. The least promising technologies are entered in the quarter called the quaking aspen with their success being either unlikely or impossible. The evaluation of the technology value made has revealed that all the groups of technologies are characterised by a high potential and attractiveness, thus classified to the most promising quarter called the

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201

Investigations of processing technology of the used alloys were tightened to the area of surface engineering only, because there exist a wide variety of available types of coatings and methods for forming structure and properties of the surface layer of engineering materials, including magnesium alloys. The investigations allow to determine precise and complete the way to design preferably summarized core material properties as well as the surface layer of the

In addition, the presented "hard" investigation results concerns the characteristics of syner‐ gistic effect of the surface heat treatment on the structure and properties of the cast of mag‐ nesium alloys, aluminium alloys and steel. Surface treatment of the analysed materials was made using laser surface treatment methods, including, in particular, laser alloying and feeding of hard ceramic particles in the surface of the material to allow formation of quasi-MMCs composite structure (Metal Matrix Composites) on the surface. Laser surface feeding and alloying was conducted by remelting of magnesium, aluminium alloys and steel surface with hard carbide and oxide particles. The feeding and alloying materials were BN, Si3N4, ZrO2, Al2O3, NbC, TaC, VC, WC, TiC, SiC powders. The powders was supplied by side injection rate of 6±1 g/min. The laser feeding was performed by high power laser diode HPDL Rofin DL 020 using the argon shielding gas. Argon was used during laser remelting to prevent oxidation of the coating and the substrate. The process parameters during the present investigation were: laser power-1.2, 1.6, 2.0 kW, scan rate – 0.5-1.0 m/min.

The observations of the investigated materials have been made on the light microscope LEICA MEF4A as well as on the scanning electron microscope (SEM) ZEISS SUPRA 35 and on the JEOL JEM 3010 transmission electron microscope (TEM), at the accelerating voltage of 300 kV.

Hardness tests were made using Zwick ZHR 4150 TK hardness tester in the HRF scale. Microhardness of the cross section of the laser surface remelted layer was measured on Fully-Automatic Microhardness Testing System with a loading time of 15 s and the testing load of 100 g.

#### **3. Investigation results**

**2. Investigation method**

200 Light Metal Alloys Applications

aluminium and magnesium alloys.

100 g.

laser power-1.2, 1.6, 2.0 kW, scan rate – 0.5-1.0 m/min.

Investigations described in this chapter concerns first of all investigation of structure and properties of the described engineering materials, in the state after heat and surface treatment, including the selected group of magnesium alloys, aluminium alloys and steel, as well as heuristic analysis using the methodology of integrated computer prediction, enabling the demarcation of predicted development trends of the analysed groups of materials and technologies of surface treatment and lead to the determination of their strategic development in comparison with other groups of materials and surface engineering technologies. In the initial planning stage of research scope in this investigations, in the context of choosing specific research topics with a comprehensive range of academic disciplines or specialties an innova‐ tive analysis was used to select the technology that uses the complex dendro-chronological evaluation matrix. Furthermore, in order to confirm the appropriateness of the proper selection of the investigated substrate material like the described in this paper Mg-Al-Zn cast magne‐ sium alloys, there was performed an analysis using the dendrological selection matrix, considering only three groups of most commonly used construction materials, namely steel,

In addition, the presented "hard" investigation results concerns the characteristics of syner‐ gistic effect of the surface heat treatment on the structure and properties of the cast of mag‐ nesium alloys, aluminium alloys and steel. Surface treatment of the analysed materials was made using laser surface treatment methods, including, in particular, laser alloying and feeding of hard ceramic particles in the surface of the material to allow formation of quasi-MMCs composite structure (Metal Matrix Composites) on the surface. Laser surface feeding and alloying was conducted by remelting of magnesium, aluminium alloys and steel surface with hard carbide and oxide particles. The feeding and alloying materials were BN, Si3N4, ZrO2, Al2O3, NbC, TaC, VC, WC, TiC, SiC powders. The powders was supplied by side injection rate of 6±1 g/min. The laser feeding was performed by high power laser diode HPDL Rofin DL 020 using the argon shielding gas. Argon was used during laser remelting to prevent oxidation of the coating and the substrate. The process parameters during the present investigation were:

The observations of the investigated materials have been made on the light microscope LEICA MEF4A as well as on the scanning electron microscope (SEM) ZEISS SUPRA 35 and on the JEOL JEM 3010 transmission electron microscope (TEM), at the accelerating voltage of 300 kV.

Hardness tests were made using Zwick ZHR 4150 TK hardness tester in the HRF scale. Microhardness of the cross section of the laser surface remelted layer was measured on Fully-Automatic Microhardness Testing System with a loading time of 15 s and the testing load of

#### **3.1. First part – Heuristic analysis**

In the first step of the heuristic material investigation there should be distinguished among all the available constructional materials and technologies several homogeneous groups of materials for the reason to undertake the planned experimental and comparative analysis. The impact of external positive factors (opportunities) and negative factors (difficulties) on the specific technologies analysed was evaluated using the metrological matrix of environment influence. In consistency with the concept proposed, each of the technologies assessed by the experts for their position against the macro-and microenvironment is placed in one of the quarters of this matrix. A universal scale of relative states, being a single-pole positive scale without zero, where 1 is a minimum value and 10 an extraordinarily high value, was used in the research undertaken. The dendrological matrix of values allows in a relatively easy way to visualize the results of the assessment of individual groups of materials and in a next steps of processing technologies in terms of potential, which is the real objective value of the studied area, the so-called hard values, tangible qualities as well as attractiveness, which reflects the subjective perception of its potential users,the so-called soft properties. The potential of a given group presented on the horizontal axis is the result of multi-criteria analysis carried out on the basis of expertise and extensive knowledge of literature data, taking into account the appro‐ priate proportions of the potential: creativeness, application, quality, and technical develop‐ ment. The vertical axis shows the importance of the attractiveness of the group, which is average mean value of expert and literature data evaluation in the field of research carried out on the basis of specific criteria corresponding to the economic attractiveness, cultural, eco‐ nomic, scientific and systemic. The technologies analysed were first evaluated by key experts for their attractiveness and potential and the result obtained was entered into one of the quarters of dendrological matrix of technology value. The wide-stretching oak is the most promising quarter guaranteeing the future success encompassing the technologies with a high potential and attractiveness. The soaring cypress includes very attractive technologies with a limited potential and rooted dwarf mountain pine symbolises technologies with a high potential and limited attractiveness likely to ensure a strong position for the technology if an adequate strategy is employed. The least promising technologies are entered in the quarter called the quaking aspen with their success being either unlikely or impossible. The evaluation of the technology value made has revealed that all the groups of technologies are characterised by a high potential and attractiveness, thus classified to the most promising quarter called the wide-stretching oak. (Fig. 1).

Investigations of processing technology of the used alloys were tightened to the area of surface engineering only, because there exist a wide variety of available types of coatings and methods for forming structure and properties of the surface layer of engineering materials, including magnesium alloys. The investigations allow to determine precise and complete the way to design preferably summarized core material properties as well as the surface layer of the produced element.

**Figure 1.** Dendrological matrix of technology value; approach presentation [by A.D. Dobrzańska-Danikiewicz]

**Figure 2.** Dendrochronology matrix of values presenting the placement of evaluated groups of materials investigated

Heuristic Analysis Application for Magnesium Alloys Properties Improvement

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203

**Figure 3.** Dendrochronology matrix of values presenting the placement of evaluated groups of technologies investi‐

using computer integrated forecasting of the development

gated using computer integrated forecasting of the development

For determination of the value of each group of materials, and in the second step of their surface treatment technology, there was applied the weighted scoring method for a comparative assessment used for appropriate qualifications for the specific groups of materials and technologies in the context of their mutual relationship. It should be noted that this technique involves the application of the relativistic principle of evaluation criteria, namely the recog‐ nition of differences and existing rules between the different research areas, assumed a specific set of criteria, acting as a selective filter evaluating positively or negatively the investigated object. The weighted scoring method enables the use of multi-criteria evaluation of areas which acts as the perfect model of synthetic evaluation of specific areas, using an interval scale.

Detailed criteria for evaluation of the attractiveness and potential of investigated areasconcerning materials and technologies-investigated in this work are presented in Tables 1 and 3. The listed criteria assigned specific importance, used for following calculation of the weighted values and their sum, which is the basis for comparative analysis, as shown in Tables 2 and 4. Then the multi-criteria results are visualized using a matrix of dendrological values for the investigated area (Figs. 2 and 3). The analysis of preferences clearly indicates, that the highest weighted evaluation, taking into account the criteria in each of the group of materials, are determined for magnesium alloys, which meet a wide range of demands and expectations of modern engineering materials imposed by the ever-changing environment, marked in Fig. 2 as M1 and coloured green. The investigations confirms also, that the aluminium alloys are characterized by a high level of attractiveness.

Heuristic Analysis Application for Magnesium Alloys Properties Improvement http://dx.doi.org/10.5772/57587 203

**Figure 2.** Dendrochronology matrix of values presenting the placement of evaluated groups of materials investigated using computer integrated forecasting of the development

For determination of the value of each group of materials, and in the second step of their surface treatment technology, there was applied the weighted scoring method for a comparative assessment used for appropriate qualifications for the specific groups of materials and technologies in the context of their mutual relationship. It should be noted that this technique involves the application of the relativistic principle of evaluation criteria, namely the recog‐ nition of differences and existing rules between the different research areas, assumed a specific set of criteria, acting as a selective filter evaluating positively or negatively the investigated object. The weighted scoring method enables the use of multi-criteria evaluation of areas which acts as the perfect model of synthetic evaluation of specific areas, using an interval scale.

**Figure 1.** Dendrological matrix of technology value; approach presentation [by A.D. Dobrzańska-Danikiewicz]

*P O T E N T I A L O F T E C H N O L O G Y*  1 2,5 4 5,5 7 8,5 10

*Technologies having huge attractiveness and potential*

*Technologies having limited attractiveness and huge potential*

> **Rooted DWARF MOUNTAIN PINE**  POSSIBLE SUCCESS

**Wide-stretching OAK**  GUARANTEED SUCCESS

Detailed criteria for evaluation of the attractiveness and potential of investigated areasconcerning materials and technologies-investigated in this work are presented in Tables 1 and 3. The listed criteria assigned specific importance, used for following calculation of the weighted values and their sum, which is the basis for comparative analysis, as shown in Tables 2 and 4. Then the multi-criteria results are visualized using a matrix of dendrological values for the investigated area (Figs. 2 and 3). The analysis of preferences clearly indicates, that the highest weighted evaluation, taking into account the criteria in each of the group of materials, are determined for magnesium alloys, which meet a wide range of demands and expectations of modern engineering materials imposed by the ever-changing environment, marked in Fig. 2 as M1 and coloured green. The investigations confirms also, that the aluminium alloys are

characterized by a high level of attractiveness.

*A*

2,5

4

5,5

7

8,5

10

*Technologies having limited potential and huge attractiveness*

*Technologies having limited attractiveness and potential*

POSSIBLE SUCCESS

**Quaking ASPEN**  LITTLE POSSIBLE SUCESS

**Soaring CYPREES**

*T*

*RA C T I V E N E S S O F TE*

*C*

*H*

*N*

*O*

*L*

*O*

*G Y*

202 Light Metal Alloys Applications

**Figure 3.** Dendrochronology matrix of values presenting the placement of evaluated groups of technologies investi‐ gated using computer integrated forecasting of the development

T3 points-selected as the most suitable surface treatment technologies for magnesium alloys, far exceeding the other analysed coated methods. Other techniques were evaluated in the area of the quadrant of trembling aspen or soaring cypress which may indicate, in the case of technology in the area of a trembling aspen a slow displacement of these methods by newer

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However, in case of technologies, which are placed in the area of soaring cypress there it is a need for further research and investments enabling them to develop. The only exception to the assumptions may be layering paint technologies that typically are located in the square of rooted dwarf mountain pine, in relation to painting. This technology is unattractive, but widely used due to its simple mechanism for layering and low application cost. However, according to the established criteria concerning the potential and attractiveness for future application specific to the substrate materials (magnesium alloy); the painting techniques are qualified for the quarter trembling aspen. The selected technologies-T1, T2 and T3 were characterised in

**No. Potential Importance**

gradient, composite, metastable, nanocrystalline) 0,3

(corrosion resistant, wear resistant) 0,2

composition and structure 0,1

reactions and the need for their recycling) 0,2

methods 0,3

**No. Attractiveness Importance**

Criterion 1 The possibility to obtain complex coatings (complex, multi-layered, multi-phase,

Criterion 3 The possibility to obtain hard surface layers with special protective properties

Criterion 1 Environmental friendly deposition process (no harmful products of chemical

Criterion 2 Possibility to produce coatings with properties impossible to obtain with other

Criterion 3 Wide range of further development of the technology 0,2

Criterion 4 Possibility of full automation of the manufacturing process of coating 0,1

Criterion 5 Necessity for high precision of the deposition processes 0,2

**Table 3.** Detailed criteria for the assessment of the potential and attractiveness of the group of investigated material

Criterion 2 Large choice of coating material; wide range of deposited coatings properties 0,2

Criterion 4 Layers with good adhesion to the substrate 0,2

Criterion 5 The possibility to obtain in a single process step graded layers with defined chemical

more environmentally friendly and versatile technology.

detail in the following description.

science and heuristic technologies

**Table 1.** Particular criteria for evaluation of the potential and attractiveness of material groups using for material engineering-heuristic investigations

**Table 2.** Multi-criterion analysis of the investigated materials groups using engineering-heuristic methods

The highest evaluation are replaced in the group of technologies (the quadrant of a large oak) by laser technologies and PVD/CVD methods, indicated in Fig. 3 as green coloured T1, T2 and T3 points-selected as the most suitable surface treatment technologies for magnesium alloys, far exceeding the other analysed coated methods. Other techniques were evaluated in the area of the quadrant of trembling aspen or soaring cypress which may indicate, in the case of technology in the area of a trembling aspen a slow displacement of these methods by newer more environmentally friendly and versatile technology.

**No. Potential Importance**

Criterion 1 Low specific gravity of the material 0,4

Criterion 2 Ability for vibration damping 0,1

Criterion 3 Good castability 0,1

Criterion 4 High mechanical properties and good weldability 0,2

Criterion 5 High corrosion resistance 0,2

**No. Attractiveness Importance**

Criterion 1 Availability of the half-finished product (raw material, chemical element) 0,4

Criterion 2 Recycling possibility 0,1

Criterion 3 Low production costs and the final price of the produced material 0,2

Criterion 4 Low level of complexity of the production technology 0,1

deformation, surface treatment)

**Criterion 1** **Criterion 2** **Criterion 3** **Criterion 4**

*M3* Steels 1,6 0,4 0,5 1,8 1,8 6,1 2 0,4 1,6 0,8 1,6 6,4

The highest evaluation are replaced in the group of technologies (the quadrant of a large oak) by laser technologies and PVD/CVD methods, indicated in Fig. 3 as green coloured T1, T2 and

**Table 2.** Multi-criterion analysis of the investigated materials groups using engineering-heuristic methods

The variety of technologies for semi-fabrication of the material and the wide choice of processing technologies possible to implement (heat treatment, plastic

**Table 1.** Particular criteria for evaluation of the potential and attractiveness of material groups using for material

**Potential Attractiveness**

**Criterion 5**

alloys 2,8 0,5 0,7 0,8 1,2 6 2,8 0,7 1,2 0,8 1,6 7,1

**weighted average**

4 0,8 0,8 0,6 0,2 **6,4** 4 0,9 1,2 0,4 1,6 **8,1**

**Criterion 1** **Criterion 2** **Criterion 3**

**Criterion 4**

**Criterion 5**

**weighted average**

0,2

Criterion 5

204 Light Metal Alloys Applications

*M1*

*M2*

engineering-heuristic investigations

**Designation Group of**

**material**

Magnesium alloys

Aluminium

However, in case of technologies, which are placed in the area of soaring cypress there it is a need for further research and investments enabling them to develop. The only exception to the assumptions may be layering paint technologies that typically are located in the square of rooted dwarf mountain pine, in relation to painting. This technology is unattractive, but widely used due to its simple mechanism for layering and low application cost. However, according to the established criteria concerning the potential and attractiveness for future application specific to the substrate materials (magnesium alloy); the painting techniques are qualified for the quarter trembling aspen. The selected technologies-T1, T2 and T3 were characterised in detail in the following description.

**Table 3.** Detailed criteria for the assessment of the potential and attractiveness of the group of investigated material science and heuristic technologies

fatigue (Fig. 5), especially for VC-alloyed steel surface (Figs. 6, 7 and 10, 11) compared to the

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**Figure 4.** Hardness measurement results of the 32CrMoV12-28 steel alloyed with ceramic powders

surface roughness compared to the VC alloying variant (Figs. 8, 9 and 12, 13).

Similar relationship can be found in case of Si3N4 powder alloying, with lower hardness and

The performed investigations of the alloyed hot work tool steel 32CrMoV12-28 show a clear effect of the applied ceramic powder used for alloying (Figs. 8, 9). It can be also clearly recognised the influence of the used laser power in the range of 1.2; 1.6; 2.0 and 2.3 kW on the shape and thickness as well as the particle distribution of the remelted material. It can be seen that with the increasing laser power the distribution of the remelted metal in the steel substrate increases. Microstructure presented on Figs. 8, 9 shows a dendritic structure in the remelted area. There are also VC particles present distributed in the matrix confirmed by the transmis‐ sion electron microscope using the selected area diffraction pattern. There is also a clear relationship between the employed laser power and the dendrite size, namely with increasing laser power the dendrites are larger, this relationship is valid for the Si3N4 powder alloying.

traditionally treated materials.

**Table 4.** The results of multi-criteria analysis of the investigated groups of technologies

#### **3.2. Second part – Material engineering investigations**

The developed technologies focusing on hybrid surface layers, so-called quasi-composite MMCs structures (characterized by phase composition gradient, chemical composition and functional gradient) in the process of laser alloying and/or feeding with ceramic particles into the surface of the treated materials provide a complete and comprehensive solution for modelling of engineering materials. Currently, the concept of functional lightweight materials is a priority and the most investigated worldwide field of material science and engineering concerning the production and processing of new developed engineering materials. Previous studies about the effects of laser beam effect on various materials, including magnesium alloys as well as tool steels, based on the authors own long-term investigation, summarizing the experience of laser surface treatment reveal that, there are chemical composition and structure changes which are different from those occurring during conventional heat treatment. This causes, that the laser treated elements shows relatively high hardness (Fig. 4) and thermal fatigue (Fig. 5), especially for VC-alloyed steel surface (Figs. 6, 7 and 10, 11) compared to the traditionally treated materials.

**Designation**

**Technology group**

*T1* PVD techniques, including also the Cathodic Arc Deposition

206 Light Metal Alloys Applications

*T2* CVD techniques,

*T3* Laser feeding/

*T11* Hybrid technologies

**Criterion 1**

**Criterion 2**

**Criterion 3**

**Criterion 4**

**Criterion 5**

including the PACVD 2,7 1,6 1,8 1,2 0,8 **8,1** 1,4 2,7 1,6 0,9 1 **7,6**

alloying/remelting 2,7 1 1,8 2 0,5 **8** 1,8 2,7 2 1 2 **9,5**

(multiplex) 1,5 1 1,2 1,2 0,5 5,4 1 2,7 1,2 0,5 0,8 6,2

The developed technologies focusing on hybrid surface layers, so-called quasi-composite MMCs structures (characterized by phase composition gradient, chemical composition and functional gradient) in the process of laser alloying and/or feeding with ceramic particles into the surface of the treated materials provide a complete and comprehensive solution for modelling of engineering materials. Currently, the concept of functional lightweight materials is a priority and the most investigated worldwide field of material science and engineering concerning the production and processing of new developed engineering materials. Previous studies about the effects of laser beam effect on various materials, including magnesium alloys as well as tool steels, based on the authors own long-term investigation, summarizing the experience of laser surface treatment reveal that, there are chemical composition and structure changes which are different from those occurring during conventional heat treatment. This causes, that the laser treated elements shows relatively high hardness (Fig. 4) and thermal

**Table 4.** The results of multi-criteria analysis of the investigated groups of technologies

**3.2. Second part – Material engineering investigations**

*T4* Thermal spraying 1,5 1 1,4 0,6 0,5 5 1,2 1,8 0,6 0,9 1,4 5,9 *T5* Anodisation 0,6 0,2 1,2 1,2 0,3 3,5 1,2 1,5 0,2 0,4 1 4,3 *T6* Galvanic technologies 0,9 1 1,2 1,4 0,2 4,7 0,4 0,9 0,2 0,9 0,4 2,8 *T7* Laser ablation-PLD 1,5 1 1 1,4 0,5 5,4 1,6 1,2 1,2 0,6 1,2 5,8 *T8* Painting coating 0,3 0,8 1 1,4 0,1 3,6 0,6 0,6 0,2 0,9 0,8 3,1 *T9* Ion implantation 1,5 1 1 1,2 0,5 5,2 1,4 1,8 1,2 0,6 1 6 *T10* Arc nitriding 0,6 0,4 1,2 1,2 0,3 3,7 0,8 1,8 0,6 0,8 0,8 4,8

**weighted average**

**Potential Attractiveness**

**Criterion 1**

2,7 1,6 1,8 1,2 0,9 **8,2** 1,8 2,7 1,6 0,9 1 **8**

**Criterion 2**

**Criterion 3**

**Criterion 4**

**Criterion 5**

**weighted average**

**Figure 4.** Hardness measurement results of the 32CrMoV12-28 steel alloyed with ceramic powders

Similar relationship can be found in case of Si3N4 powder alloying, with lower hardness and surface roughness compared to the VC alloying variant (Figs. 8, 9 and 12, 13).

The performed investigations of the alloyed hot work tool steel 32CrMoV12-28 show a clear effect of the applied ceramic powder used for alloying (Figs. 8, 9). It can be also clearly recognised the influence of the used laser power in the range of 1.2; 1.6; 2.0 and 2.3 kW on the shape and thickness as well as the particle distribution of the remelted material. It can be seen that with the increasing laser power the distribution of the remelted metal in the steel substrate increases. Microstructure presented on Figs. 8, 9 shows a dendritic structure in the remelted area. There are also VC particles present distributed in the matrix confirmed by the transmis‐ sion electron microscope using the selected area diffraction pattern. There is also a clear relationship between the employed laser power and the dendrite size, namely with increasing laser power the dendrites are larger, this relationship is valid for the Si3N4 powder alloying.

In general the hot work tool steel has a ferritic structure with homogeny distributed carbides in the metal matrix in the annealed state. In areas, which are between the solid and molten state dendritic structure with large dendrites can be found in case of Si3N4 powder (Fig. 14) or a fine grained structure in case of VC alloying (Fig. 15). The EDS point wise chemical compo‐ sition analysis confirms the presence of carbide ceramic particles in the matrix in form of big conglomerates. The required hardenability for this tool steel was achieving after a suitable tempering time, which assures melting of the alloying carbides in the austenite. The structural investigations carried out using the high power diode laser allows to compare the surface layer as well as the shape and depth of the remelting area. It was noticed that the depth of remelting area grows together with the increasing laser power (Fig. 16).

**Figure 6.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

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**Figure 7.** 32CrMoV12-28 steel alloyed with VC powder, laser power 2.3 kW

**Figure 8.** .32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

**Figure 5.** Influence of laser power and the type of alloyed surface layer of the 32CrMoV12-28 steel on mean crack depth, maximal crack depth and crack density occurred during the thermal fatigue test

It was found, that in case of VC powder the difference of the remelted area thickness is several times higher in case of 2.3 kW power compared to 1.2 kW laser power. Also for the same laser power (2.3 kW) the surface layer thickness increases from 1.9 mm for WC powder to 2.2 mm for TiC powder and to 2.3 mm in case of the VC powder (Fig. 16).

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**Figure 6.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

In general the hot work tool steel has a ferritic structure with homogeny distributed carbides in the metal matrix in the annealed state. In areas, which are between the solid and molten state dendritic structure with large dendrites can be found in case of Si3N4 powder (Fig. 14) or a fine grained structure in case of VC alloying (Fig. 15). The EDS point wise chemical compo‐ sition analysis confirms the presence of carbide ceramic particles in the matrix in form of big conglomerates. The required hardenability for this tool steel was achieving after a suitable tempering time, which assures melting of the alloying carbides in the austenite. The structural investigations carried out using the high power diode laser allows to compare the surface layer as well as the shape and depth of the remelting area. It was noticed that the depth of remelting

**Figure 5.** Influence of laser power and the type of alloyed surface layer of the 32CrMoV12-28 steel on mean crack

It was found, that in case of VC powder the difference of the remelted area thickness is several times higher in case of 2.3 kW power compared to 1.2 kW laser power. Also for the same laser power (2.3 kW) the surface layer thickness increases from 1.9 mm for WC powder to 2.2 mm

depth, maximal crack depth and crack density occurred during the thermal fatigue test

for TiC powder and to 2.3 mm in case of the VC powder (Fig. 16).

area grows together with the increasing laser power (Fig. 16).

208 Light Metal Alloys Applications

**Figure 7.** 32CrMoV12-28 steel alloyed with VC powder, laser power 2.3 kW

**Figure 8.** .32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

**Figure 9.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 2.3 kW

**Figure 12.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

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**Figure 13.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 2.3 kW

**Figure 14.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

**Figure 10.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

**Figure 11.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

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**Figure 12.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

**Figure 9.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 2.3 kW

210 Light Metal Alloys Applications

**Figure 10.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

**Figure 11.** 32CrMoV12-28 steel alloyed with VC powder, laser power 1.2 kW

**Figure 13.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 2.3 kW

**Figure 14.** 32CrMoV12-28 steel alloyed with Si3N4 powder, laser power 1.2 kW

**Figure 15.** 32CrMoV12-28 steel alloyed with VC powder, 1.2 kW

**Figure 17.** Thin foil structure of the 32CrMoV12-28 steel after VC alloying, laser power– 2,0 kW, a) bright field, b) dif‐

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Similar relationships were found in case of magnesium alloys, where the used WC, VC, SiC, Al2O3 and TiC powder particles are present in the laser treated magnesium surface. Effect of laser feeding conditions, namely: laser power, feeding speed, type of the used ceramic powder and the applied substrate on hardness and hardness increase of the surface layer of the cast magnesium samples were investigated using the Rockwell hardness method. The measured hardness of the surface was obtained in the range from 32.4 to 105.1 HRF (Fig. 18). As a result of the performed investigations it was found, that the highest hardness increase was observed in case of MCMgAl6Znl and MCMgAl3Znl magnesium cast alloys – as materials with low concentration (<6%) of aluminium, which is mainly responsible for precipitation strengthening

Also in this case the surface layer is relatively rough, compared to the non-treated material (Fig. 19, 20). After laser feeding, there was revealed-based on the performed metallographic investigations which were carried out on light microscope-the presence of several zones in the remelted surface layer of the cast magnesium alloys, with the thickness and the powder

fraction pattern of the VC particle revealed in Fig. 17a, c) solution of the diffraction pattern from Fig. 17c

of the studied, laser treated alloys, fed with ceramic particles.

**Figure 16.** Influence of laser power on thickness of the remelted zone RZ, heat affected zone HAZ and the surface layer 32CrMoV12-28 steel after laser alloyed.

**Figure 15.** 32CrMoV12-28 steel alloyed with VC powder, 1.2 kW

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layer 32CrMoV12-28 steel after laser alloyed.

**Figure 16.** Influence of laser power on thickness of the remelted zone RZ, heat affected zone HAZ and the surface

**Figure 17.** Thin foil structure of the 32CrMoV12-28 steel after VC alloying, laser power– 2,0 kW, a) bright field, b) dif‐ fraction pattern of the VC particle revealed in Fig. 17a, c) solution of the diffraction pattern from Fig. 17c

Similar relationships were found in case of magnesium alloys, where the used WC, VC, SiC, Al2O3 and TiC powder particles are present in the laser treated magnesium surface. Effect of laser feeding conditions, namely: laser power, feeding speed, type of the used ceramic powder and the applied substrate on hardness and hardness increase of the surface layer of the cast magnesium samples were investigated using the Rockwell hardness method. The measured hardness of the surface was obtained in the range from 32.4 to 105.1 HRF (Fig. 18). As a result of the performed investigations it was found, that the highest hardness increase was observed in case of MCMgAl6Znl and MCMgAl3Znl magnesium cast alloys – as materials with low concentration (<6%) of aluminium, which is mainly responsible for precipitation strengthening of the studied, laser treated alloys, fed with ceramic particles.

Also in this case the surface layer is relatively rough, compared to the non-treated material (Fig. 19, 20). After laser feeding, there was revealed-based on the performed metallographic investigations which were carried out on light microscope-the presence of several zones in the remelted surface layer of the cast magnesium alloys, with the thickness and the powder particles occurrence depending on the laser processing parameters and the used powder and substrate type (Figs. 21, 22). Starting from the top zone of the surface layer there occurs one zone rich in non-dissolved particles located on the surface of magnesium alloys, the second zone is the remelting area zone (RZ), with the thickness and shape directly depending on the applied laser power. Finally there occurs also the heat affected zone (HAZ) (Figs. 23, 24).

**Figure 20.** Surface layer of the MCMgAl3Zn1 alloy after SiC powder feeding, laser power 1.6 kW, laser scan rate 0.75

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**Figure 21.** Surface layer of the MCMgAl3Zn1 alloy after WC powder feeding, laser power 1.6 kW, laser scan rate 0.75

**Figure 22.** Surface layer of the MCMgAl9Zn1 alloy after WC powder feeding, laser power 2.0 kW, laser scan rate 0.75

m/min

m/min

m/min

**Figure 18.** Hardness measurements results of cast Mg-Al-Zn magnesium alloys samples, after aging and laser feeding

**Figure 19.** Surface layer of the MCMgAl9Zn1 alloy after WC powder feeding, laser power 2.0 kW, laser scan rate 0.5 m/min

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particles occurrence depending on the laser processing parameters and the used powder and substrate type (Figs. 21, 22). Starting from the top zone of the surface layer there occurs one zone rich in non-dissolved particles located on the surface of magnesium alloys, the second zone is the remelting area zone (RZ), with the thickness and shape directly depending on the applied laser power. Finally there occurs also the heat affected zone (HAZ) (Figs. 23, 24).

**Figure 18.** Hardness measurements results of cast Mg-Al-Zn magnesium alloys samples, after aging and laser feeding

**Figure 19.** Surface layer of the MCMgAl9Zn1 alloy after WC powder feeding, laser power 2.0 kW, laser scan rate 0.5

m/min

214 Light Metal Alloys Applications

**Figure 20.** Surface layer of the MCMgAl3Zn1 alloy after SiC powder feeding, laser power 1.6 kW, laser scan rate 0.75 m/min

**Figure 21.** Surface layer of the MCMgAl3Zn1 alloy after WC powder feeding, laser power 1.6 kW, laser scan rate 0.75 m/min

**Figure 22.** Surface layer of the MCMgAl9Zn1 alloy after WC powder feeding, laser power 2.0 kW, laser scan rate 0.75 m/min

**Figure 23.** Central zone of the surface layer of the MCMgAl6Zn1 alloy after Al2O3 powder alloying, laser power 2.0 kW, laser scan rate 0.50 m/min

**Figure 24.** Central zone of the surface layer of the MCMgAl6Zn1 alloy after SiC powder alloying, laser power 2.0 kW, laser scan rate 0.50 m/min

These zones, depending on the used laser power and the applied ceramic powder are of varying thickness and shape (Fig. 25). It is possible to obtain such an effect of significant refinement of the grains only because of a fast heat transport from the remelting liquid metal lake through magnesium substrate of high thermal capacity and very good thermal conduc‐ tivity, which results with the increase of grain boundaries amount representing a solid barrier for the dislocations movement and therefore reinforcement of the material.

The structure of the solidified material after laser treatment is characterised, like mentioned before, with a zone like structure with diversified morphology related to the crystallisation of magnesium alloys, containing particle of the alloying powders (Figs. 26, 27), serving also like dislocation barrier. Multiple change of crystal growth direction has been observed for these areas. In the area located on the boundary between the solid and liquid phase, minor dendrites occur with their main axes oriented along to the heat transport directions.

**Figure 26.** Structure of the MCMgAl6Zn1 magnesium cast alloy after laser feeding with WC powder, laser power 1.6 kW, a) bright field, b) dark field image from the [0-13] reflection c) diffraction pattern from the particle in Fig. a, d)

**Figure 25.** Influence of laser power on thickness of the remelted zone RZ, heat affected zone HAZ and the surface

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solution of the diffraction pattern from Fig. c

layer of cast magnesium alloys after laser feeding

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**Figure 23.** Central zone of the surface layer of the MCMgAl6Zn1 alloy after Al2O3 powder alloying, laser power 2.0

**Figure 24.** Central zone of the surface layer of the MCMgAl6Zn1 alloy after SiC powder alloying, laser power 2.0 kW,

These zones, depending on the used laser power and the applied ceramic powder are of varying thickness and shape (Fig. 25). It is possible to obtain such an effect of significant refinement of the grains only because of a fast heat transport from the remelting liquid metal lake through magnesium substrate of high thermal capacity and very good thermal conduc‐ tivity, which results with the increase of grain boundaries amount representing a solid barrier

The structure of the solidified material after laser treatment is characterised, like mentioned before, with a zone like structure with diversified morphology related to the crystallisation of magnesium alloys, containing particle of the alloying powders (Figs. 26, 27), serving also like dislocation barrier. Multiple change of crystal growth direction has been observed for these areas. In the area located on the boundary between the solid and liquid phase, minor dendrites

for the dislocations movement and therefore reinforcement of the material.

occur with their main axes oriented along to the heat transport directions.

kW, laser scan rate 0.50 m/min

216 Light Metal Alloys Applications

laser scan rate 0.50 m/min

**Figure 25.** Influence of laser power on thickness of the remelted zone RZ, heat affected zone HAZ and the surface layer of cast magnesium alloys after laser feeding

**Figure 26.** Structure of the MCMgAl6Zn1 magnesium cast alloy after laser feeding with WC powder, laser power 1.6 kW, a) bright field, b) dark field image from the [0-13] reflection c) diffraction pattern from the particle in Fig. a, d) solution of the diffraction pattern from Fig. c

The obtained results carried out by the team of authors concerning laser surface treatment of aluminium alloys reveal, that the highest hardness occurs in case of ZrO2 powder alloying. In the initial planning step of investigations, there were choose specific research topics, using the innovative science-heuristic analysis for selection of the investigated material (substrate) like the Al-Si-Cu alloys based on the hardness results (Fig. 28).

product of the heat transfer processes and may be neutralised by properly adjusted powder quality and powder feed rate. However in some cases the surface can reveal a very irregular shape (Fig. 35) with decomposition of the fed particles (Fig. 36). It was also found that the examined layers consists of three subzones – the remelted zone, the heat influence zone with

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The thickens of the powder feed depth can be determined in the range up to 2.1 mm (Fig. 39) in case of the HAZ of WC powder fed with laser power of 1.5 kW for the AlSi9Cu4 alloy

**Figure 28.** Hardness measurements results of cast Al-Si-Cu aluminium alloys samples, after aging and laser feeding

**Figure 29.** AlSi9Cu alloy Al2O3, laser power 2.0 kW, 1g/min, laser scan rate 0.5 m/min

a dendritic structure and the substrate material (Figs. 37, 38).

and 1.9 mm for the AlSi9Cu alloy.

Microstructure investigations of the aluminium alloys fed with WC, SiC Al2O3 and ZrO2 ceramic powders (Figs. 29-38) were performed light microscopy. Like in the previous case also here the roughness of the surface increases (Figs. 29, 30). It was also found out, that in case of ZrO2 powder alloying (Fig. 33) there was obtained a layer film, whereas in case of WC (Fig. 34) there are present particles in the surface layer; in both cases there are no pores or cracks in the produced coating or any defects and failures occurs in this layer. For the WC powder the particles are partially present on the bottom of the remelted zone.

**Figure 27.** Structure of the MCMgAl6Zn1 magnesium cast alloy after laser feeding with WC powder, laser power 2.0 kW, a) bright field, b) diffraction pattern from the particle in Fig. a, c) solution of the diffraction pattern from Fig. b

Further investigations should revealed the exact morphology and nature of these sublayers after alloying with different ceramic powders, carried out under different process parameters. Moreover there can be recognised, that the obtained surface are characterised by a well formed structure without any breaks or defects, they are uniformly horizontally deposited in the substrate surface. Occasionally occurred discontinuities of the layer can be recognised as a product of the heat transfer processes and may be neutralised by properly adjusted powder quality and powder feed rate. However in some cases the surface can reveal a very irregular shape (Fig. 35) with decomposition of the fed particles (Fig. 36). It was also found that the examined layers consists of three subzones – the remelted zone, the heat influence zone with a dendritic structure and the substrate material (Figs. 37, 38).

The obtained results carried out by the team of authors concerning laser surface treatment of aluminium alloys reveal, that the highest hardness occurs in case of ZrO2 powder alloying. In the initial planning step of investigations, there were choose specific research topics, using the innovative science-heuristic analysis for selection of the investigated material (substrate) like

Microstructure investigations of the aluminium alloys fed with WC, SiC Al2O3 and ZrO2 ceramic powders (Figs. 29-38) were performed light microscopy. Like in the previous case also here the roughness of the surface increases (Figs. 29, 30). It was also found out, that in case of ZrO2 powder alloying (Fig. 33) there was obtained a layer film, whereas in case of WC (Fig. 34) there are present particles in the surface layer; in both cases there are no pores or cracks in the produced coating or any defects and failures occurs in this layer. For the WC powder the

**Figure 27.** Structure of the MCMgAl6Zn1 magnesium cast alloy after laser feeding with WC powder, laser power 2.0 kW, a) bright field, b) diffraction pattern from the particle in Fig. a, c) solution of the diffraction pattern from Fig. b

Further investigations should revealed the exact morphology and nature of these sublayers after alloying with different ceramic powders, carried out under different process parameters. Moreover there can be recognised, that the obtained surface are characterised by a well formed structure without any breaks or defects, they are uniformly horizontally deposited in the substrate surface. Occasionally occurred discontinuities of the layer can be recognised as a

the Al-Si-Cu alloys based on the hardness results (Fig. 28).

218 Light Metal Alloys Applications

particles are partially present on the bottom of the remelted zone.

The thickens of the powder feed depth can be determined in the range up to 2.1 mm (Fig. 39) in case of the HAZ of WC powder fed with laser power of 1.5 kW for the AlSi9Cu4 alloy and 1.9 mm for the AlSi9Cu alloy.

**Figure 28.** Hardness measurements results of cast Al-Si-Cu aluminium alloys samples, after aging and laser feeding

**Figure 29.** AlSi9Cu alloy Al2O3, laser power 2.0 kW, 1g/min, laser scan rate 0.5 m/min

**Figure 30.** Surface layer of the AlSi9Cu alloyed with WC powder AlSi9Cu, laser power 1.5 kW, 1.5g/min, laser scan rate 0.25 m/min

**Figure 33.** Surface layer of the AlSi9Cu alloy after Al2O3 powder feeding, laser power 1.5 kW, 1g/min, laser scan rate

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**Figure 34.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 2.0 kW, 1.5g/min, laser scan rate

0.5 m/min

0.25 m/min

**Figure 31.** Cross-section of the AlSi9Cu alloy after Al2O3, powder feeding, laser power 2.0 kW, 1g/min, laser scan rate 0.5 m/min

**Figure 32.** Cross-section of the AlSi9Cu alloy, laser power 1.5 kW, 1.5g/min, laser scan rate 0.25 m/min

**Figure 30.** Surface layer of the AlSi9Cu alloyed with WC powder AlSi9Cu, laser power 1.5 kW, 1.5g/min, laser scan

**Figure 31.** Cross-section of the AlSi9Cu alloy after Al2O3, powder feeding, laser power 2.0 kW, 1g/min, laser scan rate

**Figure 32.** Cross-section of the AlSi9Cu alloy, laser power 1.5 kW, 1.5g/min, laser scan rate 0.25 m/min

rate 0.25 m/min

220 Light Metal Alloys Applications

0.5 m/min

**Figure 33.** Surface layer of the AlSi9Cu alloy after Al2O3 powder feeding, laser power 1.5 kW, 1g/min, laser scan rate 0.5 m/min

**Figure 34.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 2.0 kW, 1.5g/min, laser scan rate 0.25 m/min

**Figure 35.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 2.0 kW, 8.0 g/min, laser scan rate 0.25 m/min

**Figure 37.** Structure of the surface layer of the AlSi9Cu4 alloy, fed with SiC powder, laser power 2.0 kW, 1.5 g/min,

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**Figure 38.** Transition zone between the surface layer and AlSi9Cu4 alloy substrate, fed with SiC powder, laser power

laser scan rate 0.25 m/min

2.0 kW, 1.5 g/min, laser scan rate 0.25 m/min

**Figure 36.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 1.5 kW, 1.5 g/min, laser scan rate 0.25 m/min

**Figure 37.** Structure of the surface layer of the AlSi9Cu4 alloy, fed with SiC powder, laser power 2.0 kW, 1.5 g/min, laser scan rate 0.25 m/min

**Figure 35.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 2.0 kW, 8.0 g/min, laser scan rate

**Figure 36.** Surface layer of the AlSi9Cu alloy after WC powder feeding, laser power 1.5 kW, 1.5 g/min, laser scan rate

0.25 m/min

222 Light Metal Alloys Applications

0.25 m/min

**Figure 38.** Transition zone between the surface layer and AlSi9Cu4 alloy substrate, fed with SiC powder, laser power 2.0 kW, 1.5 g/min, laser scan rate 0.25 m/min

The metallographic examinations carried out give grounds to state that the ceramic powder alloying or feeding process will be carried out successfully in case of the aluminium alloy substrate, the powder particles will be distributed uniformly in the investigated surface layer, and that the particular layers is without cracks and failures and tightly adhere to the cast aluminium material matrix. In this part of research following conclusion can be made:

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**•** the laser surface treatment process gives satisfied results and allows to obtain high quality

**•** the surface layers consists of three zones: the remelting zone the heat influence zone and

Steel, aluminium and magnesium are the most commonly used constructional materials after laser surface treatment, moreover there are suitable in applications as components in the automotive industry. The expected range of applications for cast magnesium alloys Mg-Al-Zn surface treatment using the high performance diode laser can be interesting especially for the automotive industry, where small product weight is required, good strength properties of

This research was financed partially within the framework of the Scientific Research Project

Faculty of Mechanical Engineering, Silesian University of Technology, Gliwice, Poland

the aluminium alloys, Surface Engineering 28/8 (2012) 598-604.

loy surface layer, Solid State Phenomena 186 (2012) 192-197.

[1] T. Tański, K. Lukaszkowicz, Structure and properties of PVD coatings deposited on

[2] T. Tański, K. Labisz, Electron microscope investigation of PVD coated aluminium al‐

**•** the surface layer are without cracks and of a maximal thickness in the range of 2 mm,

substrate material as well as sometime of an additional intermediate zone.

surface in most cases of the applied ceramic powders,

components and ability to repair finished parts.

No. 2011/01/B/ST8/06663 headed by Dr Krzysztof Labisz.

Tomasz Tański, Krzysztof Labisz and Anna Dobrzańska-Danikiewicz

**Acknowledgements**

**Author details**

**References**

**Figure 39.** Influence of laser power on thickness of the remelted zone RZ, heat affected zone HAZ and the surface layer of cast aluminium alloys after laser feeding

#### **4. Investigation results**