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Effect of deep cryogenic treatment on microstructure, creep and wear behaviors of AZ91 magnesium alloy
Materials Science and Engineering A 523 (2009) 27–31
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of deep cryogenic treatment on microstructure, creep and wear behaviors of AZ91 magnesium alloy Kaveh Meshinchi Asl a,∗ , Alireza Tari b , Farzad Khomamizadeh b a b
School of Materials Science and Engineering, Clemson University, Clemson, SC, 29634, USA Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 7 September 2008 Received in revised form 30 April 2009 Accepted 4 June 2009
Keywords: Magnesium alloy Deep cryogenic treatment (DCT) Microstructure Mechanical properties Creep
a b s t r a c t This paper focuses on the effect of deep cryogenic treatment (−196 ◦ C) on microstructure and mechanical properties of AZ91 magnesium alloy. The execution of deep cryogenic treatment on samples changed the distribution of  precipitates. The tiny laminar  particles almost dissolved in the microstructure and the coarse divorced eutectic  phase penetrated into the matrix. This microstructural modification resulted in a significant improvement on mechanical properties of the alloy. The steady state creep rates were measured and it was found that the creep behavior of the alloy, which is dependent on the stability of the near grain boundary microstructure, was improved by the deep cryogenic treatment. For the AZ91 alloy, the results indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. However for the deep cryogenic samples dislocation climb controlled creep is the dominant deformation mechanism. After the deep cryogenic treatment the sliding of grain boundaries was greatly suppressed due to morphological changes. As a result, the grain boundaries are less susceptible for grain boundary sliding at high temperatures. Dry sliding wear tests were also applied and the wear resistance of the alloy improved remarkably after deep cryogenic treatment. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Cold treating is a very old process and is widely used for high precision parts. In order to avoid confusion, a distinction will be drawn between “cold treatment (CT)”, at temperatures down to about −80 ◦ C or thereabouts, and “deep cryogenic treatment (DCT)”, at about liquid–nitrogen (−196 ◦ C) temperatures. Deep cryogenic treatment improves certain properties beyond the improvement obtained by normal cold treatment. From time to time over the last few decades, interest has been shown in the effect of low temperatures during the heat treatment of steels, particularly tool steels. Deep cryogenic treatment has been shown to result in significant increases in the wear resistance of steels such as D-2 and M-2. At this time it was theorized that the increase in wear resistance was a direct result of the reduction in the amount of retained austenite and change in the carbide morphologies, however more investigations are underway to understand the mechanisms of beneficial effects of deep cryogenic treatment. The beneficial effect of cryogenics treatment on improved performance of a majority of metallic and non-metallic materials has not been presented until now. However, the special conditions which are indispens-
∗ Corresponding author. Tel.: +1 864 506 5180; fax: +1 864 656 1453. E-mail address: kaveh [email protected] (K.M. Asl). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.06.003
able for obtaining beneficial results from cryogenics treatment of materials, as well as homogenization and stabilization of internal microstructure may be the most effective reason of this observable fact occurrence. A similar approach has been used for non-ferrous alloys [1–8]. The increasing demand for reducing weight for aerospace and automotive applications has been the motivation for many research programs concerning magnesium and its alloys. Magnesium alloys offer lightweight alternatives to conventional metallic alloys because of their low density. In certain applications, lightweight alloys are subjected to sliding motion, thus wear resistance is becoming a key factor in these alloys. For critical automobile applications, wear properties of magnesium alloys are important. Despite the attractive range of bulk mechanical properties, a relatively poor resistance to fracture and wear is a serious hindrance against wider application of Mg alloys. The wear properties can be varied substantially through changes in the microstructure and the morphology or volume fraction, mechanical properties and the nature of the interface between matrix and the reinforcing phase. The most commonly used magnesium alloy, being used in approximately 90% of all magnesium cast products, is AZ91 which contains 9.0 wt.% Al and 0.9 wt.% Zn. While this alloy has good castability and exhibits adequate tensile yield strengths at ambient temperature, it faces strong challenges such as creep and wear resistance [9–12]. Thus, in this investigation deep cryogenic treatment was applied
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K.M. Asl et al. / Materials Science and Engineering A 523 (2009) 27–31
Table 1 Chemical compositions of the samples investigated. Alloy code
% Al
% Zn
% Mn
% RE
% Mg
AZ91
9.10
0.90
0.25
–
Rem.
to this alloy to study its effect on mechanical properties and creep resistance of AZ91 alloy. 2. Experimental procedure Commercial pure magnesium, aluminum and zinc (>99.9%) were used to prepare AZ91 alloy. Manganese was added as Al–25%Mn master alloy in order to remove the iron atoms in the raw materials. Alloy was melted in a mild steel crucible and during the melting process, surface of the charge and the melt were protected by Magrex 36 flux (Foseco Co.). Finally after complete defluxing, the melt was poured at the pouring temperature of 750 ◦ C into a permanent mold which was preheated up to 300 ◦ C. The chemical composition of the samples was determined by wet chemical analysis as shown in Table 1. Mechanical properties of round tensile samples (ASTMB557M standard) were measured using an Instron 4400 machine (1 mm/min strain rate). The hardness was measured using Instron Brinell hardness tester. The averages of three measurements were reported for each experiment. Creep tests were performed in air using a Santam creep test machine (ASTM-E139 standard). The displacement was measured using a strain gauge in the experiments. By plotting the displacement at various times, the strain curve was obtained and the steady state strain rate was calculated. The microstructure and phase distribution were characterized by Philips XL 40 scanning electron microscope (SEM). The elemental contents of various phases in the polished samples were determined by the energy dispersive X-ray spectroscopy (EDS) system of SEM. Cryogenic treatment of samples has been performed by placing specimens in an isolated alumina chamber which was designed by means of the heat transfer equations to estimate the thermal gradient of the chamber. The top of the chamber was covered by insulator wool after placing the samples in the chamber. This chamber was progressively immersed in a liquid nitrogen reservoir by means of an electric motor. The sample temperature was monitored by a K type thermocouple which was used to operate a step motor which lowered the sample and maintained a temperature decline at the rate of 0.5 ◦ C min−1 . Steps were about 1 mm and about 8 h was taken to reach to about −196 ◦ C. This painstaking method eliminates the probability of thermal shock and micro cracking. Specimens were held at −196 ◦ C for 20 h and then slowly brought up to approximately +25 ◦ C. Dry sliding wear test without lubricant according to ASTM G99 was conducted using a pin-on-disk apparatus. Pin specimens of diameter 5 mm and length 15 mm were prepared by rods grinding up to 1200-grit with silicon carbide, polishing with 0.05 m diamond paste and alcohol, followed by cleaning with ultrasonic equipment in pure acetone. M35 hardened tool steel of 160 mm diameter and 65 HRc hardness was used as a counterface. After each test the disc surface was ground against 1200-grit SiC paper and cleaned with acetone. The experiments were performed under 10, 20, 30, 50 and 100 N stationary normal loads and sliding speeds of 0.25, 0.5 and 1.0 m/s at various sliding distances from 250 to 2000 m. Weight loss values were determined from weight differences before and after the tests using a precise electronic balance with an accuracy of ±0.1 mg. Weight loss versus sliding distance curves were plotted and the wear rates were calculated.
Fig. 1. Micrographs of as-cast AZ91 alloy, (a) SEM micrograph of as-cast structure, (b) SEM micrograph showing two kinds of  particles, divorced eutectic and tiny lamellar particles.
The worn pins surfaces were examined using a scanning electron microscopy (SEM) and energy dispersive X-Ray spectroscopy (EDAX) to verify dominant wear mechanisms. For all the samples tested in this investigation, the samples were cut from the same place of the cast ingot and had the same microstructure in terms of grain size and precipitate distribution. 3. Results and discussion Fig. 1a shows the microstructure of as-cast AZ91 alloy which consists of ␣–Mg matrix and  (Mg17 Al12 ) phase. In the as-cast AZ91 alloy, the  phase exhibited two morphologies. The first one which is coarse with irregular morphology is eutectic . The second was tiny laminar shaped and surrounding the first one which is precipitate of  due to decreasing of solubility of aluminum with decreasing temperature (Fig. 1b). According to the Mg–Al phase diagram and the previous investigations [12,13], these two particles have the same composition. In the previous work, it was proved that although the  phase has the main strengthening effect on the Mg–Al based alloys at room temperature, it has a low melting temperature and is the main reason for poor mechanical properties of these alloys at elevated temperatures [14]. Non-equilibrium solidification is the reason that  particles exhibited two different kinds of morphologies [13]. According to the binary Mg–Al phase diagram, the composition of AZ91 alloy is close to the hypoeutectic Mg–Al alloys. Since the mold for casting in this investigation was from mild steel, the cooling rate of the alloys was high, hence, it was easy to produce divorced eutectic. After
K.M. Asl et al. / Materials Science and Engineering A 523 (2009) 27–31
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Fig. 3. Energy dispersive X-ray analysis of the eutectic  phase penetrated into the matrix.
Fig. 2. Micrographs of cryogenic treated AZ91 alloy, (a) SEM micrograph of as-cast cryogenic treated structure, (b) SEM micrograph showing changes in morphologies of two kinds of  particles.
alloy solidification, the precipitation of  occurred during continuous cooling and caused the formation of the second kind of , the tiny laminar particles. After deep cryogenic treatment was applied to as-cast specimens, the microstructure of AZ91 alloy went under some changes. The morphology of  phases in the microstructure changed impressively after cryogenic treatment as shown in Fig. 2 (The same sample was used in Figs. 1 and 2 by taking micrographs of as-cast sample, then cryogenically treating the sample and taking some micrographs again). The coarse divorced eutectic  phase penetrated the matrix (Fig. 3). As a result of this change in the microstructure, the mechanical properties of cryogenic treated samples increased compared with the as-cast specimens with no treatments as shown in Table 2. This improvement was attributed to the strengthening of the matrix against propagation of the existing defects which is due to the important role of  precipitates in the microstructure which are the main strengthening effect at room temperatures. Since the  precipitates are mainly distributed at grain boundaries, the morphology of the  particles after cryogenic treatment and the stabilization of internal microstructure are both key factors in strengthening these alloys. The creep tests were carried out to investigate the effect of cryogenic treatment on creep properties of the alloys. Creep tests were
Element
At. %
Magnesium Aluminum
60.14 39.86
performed on specimens over the temperature range of 135–200 ◦ C and at an applied stress of 48–96 MPa. By measuring the slope of the typical creep curves, the steady creep rate of the alloys can be calculated and it is shown in Fig. 4 for different alloys at 200 ◦ C. The most important result caused by cryogenic treatment was the improvement in creep resistance of the alloy. The creep resistance of AZ91 alloy was relatively poor but it was observed that cryogenic treatment resulted in remarkable improvement of creep resistance of the alloy. The conventional power law equation relating the minimum creep rate, ε˙ min , to the applied stress is: ε˙ min = A
G
n exp
−Q RT
Both n and Q are parameters of the material and together may be used to identify the dominant creep mechanism for the material. By plotting logarithmically the minimum creep strain rate versus the applied stress , we can calculate the stress exponent, n. Plot of log ε versus 1/T will yield the apparent activation energy, Q [15–18]. The stress exponent n ∼2 is generally reported for grain boundary sliding, while n = 3–7 is for dislocation climb controlled creep,
Table 2 Ultimate tensile strength and hardness of samples. Sample
Yield strength (MPa)
Ultimate tensile strength (MPa)
Hardness (BHN)
As-cast After cryogenic
91 98
170 187
59 67
Fig. 4. Steady state creep rates of the alloys for different stresses in 200 ◦ C.
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K.M. Asl et al. / Materials Science and Engineering A 523 (2009) 27–31
Fig. 5. Variation of creep rates with stress for alloys tested at 200 ◦ C.
however high n values for high stresses often indicate of power low breakdown [18,19]. The activation energy for self-diffusion in magnesium can be taken as 138 kJ mol−1 . The value of activation energy for boundary diffusion is estimated to be in the range 70–100 kJ mol−1 [18]. At low stresses, the stress exponent for as-cast AZ91 was 5.6 which is in the range of dislocation climb mechanism however the activation energy was measured 121 kJ mol−1 which lies between that for self-diffusion and grain boundary diffusion. Together these two factors may indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. For the cryogenic treated AZ91 alloy, the stress exponent was 3.4 and the activation energy was measured 132 kJ mol−1 . The values of both parameters fit the values for dislocation climb controlled creep (Figs. 5 and 6). The activation energy for cryogenic treated alloy is slightly bigger than AZ91 alloys which is due to its better creep resistance. The discontinues precipitation of  dominant at grain boundaries weakens the grain boundaries at elevated temperatures. Although they are the main strengthening factor of the alloy at room temperature, they can readily soften and coarsen with increase of temperature. Thus they can have detritus effects on the high temperature properties of the alloy. Therefore sliding of grain boundaries is an important factor in deformation of Mg–Al alloys at elevated temperatures. After the cryogenic treatment, the microstructure of the alloy was modified. The tiny laminar  particles were almost dissolved in the microstructure and the coarse divorced eutectic  phase penetrated the matrix. Thus suppressing the grain boundary sliding to great extent. As a result of this change in the microstructure, the creep properties of cryogenic
Fig. 6. Variation of creep rates with reciprocal temperature for alloys tested at 48 MPa.
Fig. 7. Wear rate values for (a) 0.25, (b) 0.5 and (c) 1.0 m/s sliding speed for cryogenic and AZ91 alloy.
treated samples increased compared with the as-cast specimens with no treatments. Fig. 7 represents the corresponding wear rate under different normal loads for various sliding speeds. At constant sliding speed, application of higher normal load increases the wear rate generally. On the other hand, under invariable conditions of load, enhancing sliding speed increases the wear rate. It should be noted that, improving sliding speeds from 0.25 to 0.5 m/s lead to amplification of wear rates however, experimenting at 1.0 m/s raises wear rates significantly more than 0.5 m/s. Observations reveal that wear rates increase abruptly between 30 and 50 N vertical loads at 0.25 and 0.5 m/s sliding speeds, but gradually at 1.0 m/s. Diagrams indicate that pure AZ91 magnesium alloy examined at 0.25 m/s exhibits the least wear rate except under 100 N normal load. In this fashion, deep cryogenic treatment had no positive effect on AZ91 wear resistance except in high loads. Tests performed at 0.5 m/s show a similar behavior unless AZ91 preserves its minor wear rate up to 50 N. It seems that deep cryogenic process influences the wear rate especially at higher loads. This fact may be related to the new morphological microstructure of  phase which has better wear resistance than ␣ phase. Indeed, deep cryogenic treatment samples when tested at 1.0 m/s display superior wear resistance compared with pure AZ91 alloy. Further investigation is in progress. Traces of parallel grooves can be distinguished on surfaces of most samples. These scratches may be attributed to hard asperities of counterface or detached particles that are removed from disk
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Deep cryogenic treatment stabilizes the internal microstructure due to new morphology of  particles phase that acquires suitable mechanical properties that enhance the load bearing restriction of AZ91 alloy. At more severe wear conditions of highest normal load and sliding speed, specimens that had undergone cryogenic treatment show superior wear resistance particularly to plastic deformation. 4. Conclusions
Fig. 8. Characteristic scheme of abrasive wear for a pure AZ91 alloy experimented at 0.5 m/s under 20 N normal load after 1000 m.
(1) Deep cryogenic treatment changes the morphology of  particles in AZ91 magnesium alloy. The coarse divorced eutectic  phase penetrated in the matrix which resulted in the improvement of mechanical properties of cryogenic treated samples compared with as-cast samples. This improvement was attributed to the strengthening of the matrix against propagation of the existing defects. (2) In AZ91 alloy the results from creep curves indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. However after deep cryogenic treatment which resulted in modifying the microstructure, dislocation climb controlled creep is the dominant creep mechanism. (3) It seems that deep cryogenic process influences the wear rate especially at higher loads. The wear resistance of deep cryogenically treated samples improved significantly at higher loads and sliding speeds which was due to the stabilization of internal microstructure and the new morphology of  particles as the main strengthening effect at room temperatures. Acknowledgment The authors acknowledge Gary Kaufman at the School of Materials Science and Engineering of Clemson University for his discussions which were very thought provoking and inspiring. References
Fig. 9. Delamination mode of wear for (a) pure AZ91 alloy tested at 0.25 m/s, 50 N and 1000 m, (b) pure AZ91 alloy tested at 0. 5 m/s, 100 N and 1000 m.
or pin and placed in surface contact [20]. Fig. 8 illustrates typical feature of such a mechanism. Abrasion is the prevailing wear mechanism at lower loads. With wear process advancement, grooves become wider and deeper at sufficiently longer sliding distances. Extensive loads at long sliding distances induce subsurface crack formation and growth. If propagated cracks joined the surface large craters will be produced [21]. Characteristic sketches of delamination were identified for long sliding distances (Fig. 9). High loads at long sliding distances, predominantly for greater sliding speeds, induce severe plastic deformation on materials [20,21].
[1] V. Firouzdor, E. Nejati, F. Khomamizadeh, Journal of Materials Processing Technology 206 (2008) 467. [2] R. Mahmudi, H.M. Ghasemi, H.R. Faradji, Heat Treatment of Metals 3 (2000) 69. [3] D.N. Collins, Heat Treatment of Metals 2 (1996) 40. [4] R.G. Bowes, Heat Treatment of Metals 1 (1974) 29. [5] D. Yun, L. Xiaoping, X. Hongshen, Heat Treatment of Metals 3 (1998) 55. [6] D.N. Collins, J. Dormer, Heat Treatment of Metals 3 (1997) 71. [7] A. Molinary, M. Pellizzari, S. Gialanella, G. Straffelini, K.H. Stiasny, Journal of Materials Processing Technology 118 (2001) 350. [8] J.Y. Huang, Y.T. Zhu, X.Z. Liao, I.J. Beyerlein, M.A. Bourke, T.E. Mitchell, Materials Science and Engineering A339 (2003) 241. [9] P.J. Blau, M. Walukas, Tribology International 33 (2000) 573. [10] H. Chen, A.T. Alpas, Wear 246 (2000) 106. [11] Y. Guangyin, S. Yangshan, D. Wenjiang, Materials Science and Engineering A308 (2001) 38. [12] D.S. Mehta, S.H. Masood, W.Q. Song, Journal of Materials Processing Technology 155–156 (2004) 1526. [13] D. Wenwen, S. Yangshan, M. Xuegang, X. Feng, Z. Min, W. Dengyun, Materials Science and Engineering A356 (2003) 1. [14] F. Khomamizadeh, B. Nami, S. Khoshkhooei, Metallurgical, Materials Transactions A (2005) 3489. [15] H.D. Kang, S.S. Park, J.N. Kim, Materials Science and Engineering A 413–414 (2005) 555. [16] S. Celotto, Acta Materialia 48 (2000) 1775. [17] D. Dully, N.C. Cheynet, Y. Bercht, Acta Metallurgica and Materialia 42 (1994) 3843. [18] C.J. Bettles, Materials Science and Engineering A348 (2003) 280. [19] H. Somekawa, K. Hirai, H. Watanabe, Y. Takigawa, K. Higashi, Materials Science and Engineering A407 (2005) 53. [20] C.Y.H. Lim, S.C. Lim, M. Gupta, Wear 255 (2003) 629. [21] S.C. Sharma, B. Anand, M. Krishma, Wear 241 (2000) 33.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of deep cryogenic treatment on microstructure, creep and wear behaviors of AZ91 magnesium alloy Kaveh Meshinchi Asl a,∗ , Alireza Tari b , Farzad Khomamizadeh b a b
School of Materials Science and Engineering, Clemson University, Clemson, SC, 29634, USA Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 7 September 2008 Received in revised form 30 April 2009 Accepted 4 June 2009
Keywords: Magnesium alloy Deep cryogenic treatment (DCT) Microstructure Mechanical properties Creep
a b s t r a c t This paper focuses on the effect of deep cryogenic treatment (−196 ◦ C) on microstructure and mechanical properties of AZ91 magnesium alloy. The execution of deep cryogenic treatment on samples changed the distribution of  precipitates. The tiny laminar  particles almost dissolved in the microstructure and the coarse divorced eutectic  phase penetrated into the matrix. This microstructural modification resulted in a significant improvement on mechanical properties of the alloy. The steady state creep rates were measured and it was found that the creep behavior of the alloy, which is dependent on the stability of the near grain boundary microstructure, was improved by the deep cryogenic treatment. For the AZ91 alloy, the results indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. However for the deep cryogenic samples dislocation climb controlled creep is the dominant deformation mechanism. After the deep cryogenic treatment the sliding of grain boundaries was greatly suppressed due to morphological changes. As a result, the grain boundaries are less susceptible for grain boundary sliding at high temperatures. Dry sliding wear tests were also applied and the wear resistance of the alloy improved remarkably after deep cryogenic treatment. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Cold treating is a very old process and is widely used for high precision parts. In order to avoid confusion, a distinction will be drawn between “cold treatment (CT)”, at temperatures down to about −80 ◦ C or thereabouts, and “deep cryogenic treatment (DCT)”, at about liquid–nitrogen (−196 ◦ C) temperatures. Deep cryogenic treatment improves certain properties beyond the improvement obtained by normal cold treatment. From time to time over the last few decades, interest has been shown in the effect of low temperatures during the heat treatment of steels, particularly tool steels. Deep cryogenic treatment has been shown to result in significant increases in the wear resistance of steels such as D-2 and M-2. At this time it was theorized that the increase in wear resistance was a direct result of the reduction in the amount of retained austenite and change in the carbide morphologies, however more investigations are underway to understand the mechanisms of beneficial effects of deep cryogenic treatment. The beneficial effect of cryogenics treatment on improved performance of a majority of metallic and non-metallic materials has not been presented until now. However, the special conditions which are indispens-
∗ Corresponding author. Tel.: +1 864 506 5180; fax: +1 864 656 1453. E-mail address: kaveh [email protected] (K.M. Asl). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.06.003
able for obtaining beneficial results from cryogenics treatment of materials, as well as homogenization and stabilization of internal microstructure may be the most effective reason of this observable fact occurrence. A similar approach has been used for non-ferrous alloys [1–8]. The increasing demand for reducing weight for aerospace and automotive applications has been the motivation for many research programs concerning magnesium and its alloys. Magnesium alloys offer lightweight alternatives to conventional metallic alloys because of their low density. In certain applications, lightweight alloys are subjected to sliding motion, thus wear resistance is becoming a key factor in these alloys. For critical automobile applications, wear properties of magnesium alloys are important. Despite the attractive range of bulk mechanical properties, a relatively poor resistance to fracture and wear is a serious hindrance against wider application of Mg alloys. The wear properties can be varied substantially through changes in the microstructure and the morphology or volume fraction, mechanical properties and the nature of the interface between matrix and the reinforcing phase. The most commonly used magnesium alloy, being used in approximately 90% of all magnesium cast products, is AZ91 which contains 9.0 wt.% Al and 0.9 wt.% Zn. While this alloy has good castability and exhibits adequate tensile yield strengths at ambient temperature, it faces strong challenges such as creep and wear resistance [9–12]. Thus, in this investigation deep cryogenic treatment was applied
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K.M. Asl et al. / Materials Science and Engineering A 523 (2009) 27–31
Table 1 Chemical compositions of the samples investigated. Alloy code
% Al
% Zn
% Mn
% RE
% Mg
AZ91
9.10
0.90
0.25
–
Rem.
to this alloy to study its effect on mechanical properties and creep resistance of AZ91 alloy. 2. Experimental procedure Commercial pure magnesium, aluminum and zinc (>99.9%) were used to prepare AZ91 alloy. Manganese was added as Al–25%Mn master alloy in order to remove the iron atoms in the raw materials. Alloy was melted in a mild steel crucible and during the melting process, surface of the charge and the melt were protected by Magrex 36 flux (Foseco Co.). Finally after complete defluxing, the melt was poured at the pouring temperature of 750 ◦ C into a permanent mold which was preheated up to 300 ◦ C. The chemical composition of the samples was determined by wet chemical analysis as shown in Table 1. Mechanical properties of round tensile samples (ASTMB557M standard) were measured using an Instron 4400 machine (1 mm/min strain rate). The hardness was measured using Instron Brinell hardness tester. The averages of three measurements were reported for each experiment. Creep tests were performed in air using a Santam creep test machine (ASTM-E139 standard). The displacement was measured using a strain gauge in the experiments. By plotting the displacement at various times, the strain curve was obtained and the steady state strain rate was calculated. The microstructure and phase distribution were characterized by Philips XL 40 scanning electron microscope (SEM). The elemental contents of various phases in the polished samples were determined by the energy dispersive X-ray spectroscopy (EDS) system of SEM. Cryogenic treatment of samples has been performed by placing specimens in an isolated alumina chamber which was designed by means of the heat transfer equations to estimate the thermal gradient of the chamber. The top of the chamber was covered by insulator wool after placing the samples in the chamber. This chamber was progressively immersed in a liquid nitrogen reservoir by means of an electric motor. The sample temperature was monitored by a K type thermocouple which was used to operate a step motor which lowered the sample and maintained a temperature decline at the rate of 0.5 ◦ C min−1 . Steps were about 1 mm and about 8 h was taken to reach to about −196 ◦ C. This painstaking method eliminates the probability of thermal shock and micro cracking. Specimens were held at −196 ◦ C for 20 h and then slowly brought up to approximately +25 ◦ C. Dry sliding wear test without lubricant according to ASTM G99 was conducted using a pin-on-disk apparatus. Pin specimens of diameter 5 mm and length 15 mm were prepared by rods grinding up to 1200-grit with silicon carbide, polishing with 0.05 m diamond paste and alcohol, followed by cleaning with ultrasonic equipment in pure acetone. M35 hardened tool steel of 160 mm diameter and 65 HRc hardness was used as a counterface. After each test the disc surface was ground against 1200-grit SiC paper and cleaned with acetone. The experiments were performed under 10, 20, 30, 50 and 100 N stationary normal loads and sliding speeds of 0.25, 0.5 and 1.0 m/s at various sliding distances from 250 to 2000 m. Weight loss values were determined from weight differences before and after the tests using a precise electronic balance with an accuracy of ±0.1 mg. Weight loss versus sliding distance curves were plotted and the wear rates were calculated.
Fig. 1. Micrographs of as-cast AZ91 alloy, (a) SEM micrograph of as-cast structure, (b) SEM micrograph showing two kinds of  particles, divorced eutectic and tiny lamellar particles.
The worn pins surfaces were examined using a scanning electron microscopy (SEM) and energy dispersive X-Ray spectroscopy (EDAX) to verify dominant wear mechanisms. For all the samples tested in this investigation, the samples were cut from the same place of the cast ingot and had the same microstructure in terms of grain size and precipitate distribution. 3. Results and discussion Fig. 1a shows the microstructure of as-cast AZ91 alloy which consists of ␣–Mg matrix and  (Mg17 Al12 ) phase. In the as-cast AZ91 alloy, the  phase exhibited two morphologies. The first one which is coarse with irregular morphology is eutectic . The second was tiny laminar shaped and surrounding the first one which is precipitate of  due to decreasing of solubility of aluminum with decreasing temperature (Fig. 1b). According to the Mg–Al phase diagram and the previous investigations [12,13], these two particles have the same composition. In the previous work, it was proved that although the  phase has the main strengthening effect on the Mg–Al based alloys at room temperature, it has a low melting temperature and is the main reason for poor mechanical properties of these alloys at elevated temperatures [14]. Non-equilibrium solidification is the reason that  particles exhibited two different kinds of morphologies [13]. According to the binary Mg–Al phase diagram, the composition of AZ91 alloy is close to the hypoeutectic Mg–Al alloys. Since the mold for casting in this investigation was from mild steel, the cooling rate of the alloys was high, hence, it was easy to produce divorced eutectic. After
K.M. Asl et al. / Materials Science and Engineering A 523 (2009) 27–31
29
Fig. 3. Energy dispersive X-ray analysis of the eutectic  phase penetrated into the matrix.
Fig. 2. Micrographs of cryogenic treated AZ91 alloy, (a) SEM micrograph of as-cast cryogenic treated structure, (b) SEM micrograph showing changes in morphologies of two kinds of  particles.
alloy solidification, the precipitation of  occurred during continuous cooling and caused the formation of the second kind of , the tiny laminar particles. After deep cryogenic treatment was applied to as-cast specimens, the microstructure of AZ91 alloy went under some changes. The morphology of  phases in the microstructure changed impressively after cryogenic treatment as shown in Fig. 2 (The same sample was used in Figs. 1 and 2 by taking micrographs of as-cast sample, then cryogenically treating the sample and taking some micrographs again). The coarse divorced eutectic  phase penetrated the matrix (Fig. 3). As a result of this change in the microstructure, the mechanical properties of cryogenic treated samples increased compared with the as-cast specimens with no treatments as shown in Table 2. This improvement was attributed to the strengthening of the matrix against propagation of the existing defects which is due to the important role of  precipitates in the microstructure which are the main strengthening effect at room temperatures. Since the  precipitates are mainly distributed at grain boundaries, the morphology of the  particles after cryogenic treatment and the stabilization of internal microstructure are both key factors in strengthening these alloys. The creep tests were carried out to investigate the effect of cryogenic treatment on creep properties of the alloys. Creep tests were
Element
At. %
Magnesium Aluminum
60.14 39.86
performed on specimens over the temperature range of 135–200 ◦ C and at an applied stress of 48–96 MPa. By measuring the slope of the typical creep curves, the steady creep rate of the alloys can be calculated and it is shown in Fig. 4 for different alloys at 200 ◦ C. The most important result caused by cryogenic treatment was the improvement in creep resistance of the alloy. The creep resistance of AZ91 alloy was relatively poor but it was observed that cryogenic treatment resulted in remarkable improvement of creep resistance of the alloy. The conventional power law equation relating the minimum creep rate, ε˙ min , to the applied stress is: ε˙ min = A
G
n exp
−Q RT
Both n and Q are parameters of the material and together may be used to identify the dominant creep mechanism for the material. By plotting logarithmically the minimum creep strain rate versus the applied stress , we can calculate the stress exponent, n. Plot of log ε versus 1/T will yield the apparent activation energy, Q [15–18]. The stress exponent n ∼2 is generally reported for grain boundary sliding, while n = 3–7 is for dislocation climb controlled creep,
Table 2 Ultimate tensile strength and hardness of samples. Sample
Yield strength (MPa)
Ultimate tensile strength (MPa)
Hardness (BHN)
As-cast After cryogenic
91 98
170 187
59 67
Fig. 4. Steady state creep rates of the alloys for different stresses in 200 ◦ C.
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Fig. 5. Variation of creep rates with stress for alloys tested at 200 ◦ C.
however high n values for high stresses often indicate of power low breakdown [18,19]. The activation energy for self-diffusion in magnesium can be taken as 138 kJ mol−1 . The value of activation energy for boundary diffusion is estimated to be in the range 70–100 kJ mol−1 [18]. At low stresses, the stress exponent for as-cast AZ91 was 5.6 which is in the range of dislocation climb mechanism however the activation energy was measured 121 kJ mol−1 which lies between that for self-diffusion and grain boundary diffusion. Together these two factors may indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. For the cryogenic treated AZ91 alloy, the stress exponent was 3.4 and the activation energy was measured 132 kJ mol−1 . The values of both parameters fit the values for dislocation climb controlled creep (Figs. 5 and 6). The activation energy for cryogenic treated alloy is slightly bigger than AZ91 alloys which is due to its better creep resistance. The discontinues precipitation of  dominant at grain boundaries weakens the grain boundaries at elevated temperatures. Although they are the main strengthening factor of the alloy at room temperature, they can readily soften and coarsen with increase of temperature. Thus they can have detritus effects on the high temperature properties of the alloy. Therefore sliding of grain boundaries is an important factor in deformation of Mg–Al alloys at elevated temperatures. After the cryogenic treatment, the microstructure of the alloy was modified. The tiny laminar  particles were almost dissolved in the microstructure and the coarse divorced eutectic  phase penetrated the matrix. Thus suppressing the grain boundary sliding to great extent. As a result of this change in the microstructure, the creep properties of cryogenic
Fig. 6. Variation of creep rates with reciprocal temperature for alloys tested at 48 MPa.
Fig. 7. Wear rate values for (a) 0.25, (b) 0.5 and (c) 1.0 m/s sliding speed for cryogenic and AZ91 alloy.
treated samples increased compared with the as-cast specimens with no treatments. Fig. 7 represents the corresponding wear rate under different normal loads for various sliding speeds. At constant sliding speed, application of higher normal load increases the wear rate generally. On the other hand, under invariable conditions of load, enhancing sliding speed increases the wear rate. It should be noted that, improving sliding speeds from 0.25 to 0.5 m/s lead to amplification of wear rates however, experimenting at 1.0 m/s raises wear rates significantly more than 0.5 m/s. Observations reveal that wear rates increase abruptly between 30 and 50 N vertical loads at 0.25 and 0.5 m/s sliding speeds, but gradually at 1.0 m/s. Diagrams indicate that pure AZ91 magnesium alloy examined at 0.25 m/s exhibits the least wear rate except under 100 N normal load. In this fashion, deep cryogenic treatment had no positive effect on AZ91 wear resistance except in high loads. Tests performed at 0.5 m/s show a similar behavior unless AZ91 preserves its minor wear rate up to 50 N. It seems that deep cryogenic process influences the wear rate especially at higher loads. This fact may be related to the new morphological microstructure of  phase which has better wear resistance than ␣ phase. Indeed, deep cryogenic treatment samples when tested at 1.0 m/s display superior wear resistance compared with pure AZ91 alloy. Further investigation is in progress. Traces of parallel grooves can be distinguished on surfaces of most samples. These scratches may be attributed to hard asperities of counterface or detached particles that are removed from disk
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Deep cryogenic treatment stabilizes the internal microstructure due to new morphology of  particles phase that acquires suitable mechanical properties that enhance the load bearing restriction of AZ91 alloy. At more severe wear conditions of highest normal load and sliding speed, specimens that had undergone cryogenic treatment show superior wear resistance particularly to plastic deformation. 4. Conclusions
Fig. 8. Characteristic scheme of abrasive wear for a pure AZ91 alloy experimented at 0.5 m/s under 20 N normal load after 1000 m.
(1) Deep cryogenic treatment changes the morphology of  particles in AZ91 magnesium alloy. The coarse divorced eutectic  phase penetrated in the matrix which resulted in the improvement of mechanical properties of cryogenic treated samples compared with as-cast samples. This improvement was attributed to the strengthening of the matrix against propagation of the existing defects. (2) In AZ91 alloy the results from creep curves indicate a mixed mode of creep behavior, with some grain boundary effects contributing to the overall behavior. However after deep cryogenic treatment which resulted in modifying the microstructure, dislocation climb controlled creep is the dominant creep mechanism. (3) It seems that deep cryogenic process influences the wear rate especially at higher loads. The wear resistance of deep cryogenically treated samples improved significantly at higher loads and sliding speeds which was due to the stabilization of internal microstructure and the new morphology of  particles as the main strengthening effect at room temperatures. Acknowledgment The authors acknowledge Gary Kaufman at the School of Materials Science and Engineering of Clemson University for his discussions which were very thought provoking and inspiring. References
Fig. 9. Delamination mode of wear for (a) pure AZ91 alloy tested at 0.25 m/s, 50 N and 1000 m, (b) pure AZ91 alloy tested at 0. 5 m/s, 100 N and 1000 m.
or pin and placed in surface contact [20]. Fig. 8 illustrates typical feature of such a mechanism. Abrasion is the prevailing wear mechanism at lower loads. With wear process advancement, grooves become wider and deeper at sufficiently longer sliding distances. Extensive loads at long sliding distances induce subsurface crack formation and growth. If propagated cracks joined the surface large craters will be produced [21]. Characteristic sketches of delamination were identified for long sliding distances (Fig. 9). High loads at long sliding distances, predominantly for greater sliding speeds, induce severe plastic deformation on materials [20,21].
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