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Structural characteristics of low-temperature plasma-nitrided layers on AISI 304 stainless steel with an α′-martensite layer
Surface and Coatings Technology 153 (2002) 125–130
Structural characteristics of low-temperature plasma-nitrided layers on AISI 304 stainless steel with an a9-martensite layer Zhiwei Yu, Xiaolei Xu*, Liang Wang, Jianbing Qiang, Zukun Hei Dalian Maritime University, Institute of Metal and Technology, Dalian 116024, PR China Received 7 March 2001; accepted in revised form 1 November 2001
Abstract The characteristics of the structure formation of a low-temperature plasma-nitrided layer on AISI 304 austenitic stainless steel with a pre-existing a9-martensitic deformation layer were studied in detail by glancing-angle X-ray diffraction (XRD). The results show a two-step formation process for the plasma-nitrided layer. The a9-martensitic layer gradually reverts to the original gaustenite structure at first, followed by the transformation of g™gN as the nitrogen content increases with treatment time. Within the same nitriding time, the sample with an a9 layer shows much less lattice expansion than the sample without an a9 layer. The volume fractions of a9 and g phases transformed during the plasma-nitriding process were calculated by quantitative XRD. The transformation of a9™g is dependent on the nitrogen content dissolved in a9, apart from that resulting from heating at the nitriding temperature. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Plasma nitriding; Stainless steel; Transformation
1. Introduction Most stainless steels (and in particular austenitic grades) have relatively low hardness, and therefore poor resistance to wear, unless a surface modification treatment can be applied. Plasma nitriding has been widely used in industrial production to increase the surface hardness, and hence improve the wear resistance of ferrous materials. Generally, the treatment temperature for traditional plasma-nitriding is above 500 8C w1x in order to obtain a sufficiently thick modified layer. However, CrN precipitation w2,3x occurs for most stainless steels when the nitriding temperature is increased above 450 8C, so that their corrosion resistance is largely decreased. Reducing the treatment temperature can avoid precipitation of CrN and a nitrogen-expanded austenite (gN) can form, but lower nitriding efficiency results, leading to increased treatment times w4x. Traditional plasma-nitriding technologies are therefore of limited use in industrial treatment of stainless steels. A new, low-pressure plasma arc-source ion nitriding technology developed by the authors w5,6x can substantially increase *Corresponding author. Tel.: q86-411-4729613; fax: q86-4114671395. E-mail address: [email protected] (X. Xu).
the efficiency of plasma-nitriding at low temperature and an adequately thick gN modified layer can be obtained in a shorter treatment time. There are many reports on gN-modified layers on austenitic stainless steels produced by plasma-nitriding, and plasma immersion (PIII) and plasma source ion implantation (PSII) w3,7x, but in most experiments, an ‘ideal’ specimen surface having a single g-phase structure was selected, which does not correspond to the typical production conditions. Usually, there is a hardened layer formed on the surface of metastable austenitic stainless steel by machining, which consists of the deformation martensite (a9) coexisting with g-austenite. There should be a difference in the structural characteristics and formation process for the nitrided layer for austenitic stainless steel with or without a pre-introduced a9 layer. In this paper, we investigate the structural characteristics of nitrided layers on 304 stainless steels with an a9 deformation layer and examine the related transformation during the plasma-nitriding process. 2. Experimental methods The composition of the annealed AISI 304 austenite stainless steel used for this work was (in wt.%): C,
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 1 . 0 1 6 7 0 - X
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Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
0.07; Si, 0.78; Mn, 0.90; S, 0.13; P, 0.03; Cr, 19.0; Ni, 8.26; and Fe, balance. Plate samples of dimensions 15=15=5 mm3 were prepared and mechanically polished using 800噛 grit SiC paper. The samples were plasma-nitrided using the above-mentioned low-pressure plasma arc source at 400–420 8C. Ammonia was used as the working gas. The working pressure was approximately 0.4 Pa. The arc current was 40 A with a voltage of 50 V. A negative bias of 0.8–1 kV was applied to the substrates. The current density on the substrate surface was 0.6–0.8 mAycm2 The phases present in the modified layer were determined by glancing-angle Xray diffraction on a Rigaku Dymax-IIIA diffractometer. The radiation used was CuKa (ls0.1542 nm) at an incident angle of 108. The X-ray penetration depth under these diffraction conditions was estimated to be approximately 2.5 mm. 3. Results and discussion 3.1. Structure of the matrix XRD patterns of the electrically and mechanically polished samples are shown in Fig. 1. The XRD pattern in Fig. 1a is typical of a f.c.c. structure, the peaks of which appear at 2us43.718, 50.928, 74.838, etc., corresponding to plane spacings of ds0.2069, 0.1792, 0.128
Table 1 Mass change of the sample during nitriding Nitriding
Mass (g)
time
Measured
Average
6.61909 6.61909 6.61909 6.61909 6.61907 6.61906 6.61907 6.61906 6.61904 6.61905 6.61904 6.61905 6.61874 6.61875 6.61875 6.61874
6.61909
0
12 min
18 min
4h
Mass loss (mg)
6.61907
0.02
6.61905
0.04
6.61875
0.34
nm, etc., respectively. This demonstrates that the electrically polished sample is composed of single-phase gaustenite. However, in the XRD pattern of Fig. 1b, the peaks (2us44.568, 64.728, 82.728) of the b.c.c. structure corresponding to the plane spacings of the ferrite (ds 0.2032, 0.1440, 0.1171 nm) appear, except for the peaks of g-austenite. Its peak intensity is higher than that of g-austenite, which demonstrates that there is a certain amount of a9-martensite induced by the strain on the surface during mechanical polishing of the sample. The peaks of the remaining g-austenite did not shift. 3.2. Structure composition of the plasma-nitrided layer
Fig. 1. XRD patterns of 304 stainless steel (a) electrically and (b) mechanically polished.
XRD patterns of samples nitrided for different times are shown in Fig. 2. It is evident that the peak intensity of a9-martensite gradually decreased, but the intensity of the g-austenite increased with treatment time. This suggests that a gradual transformation of a9 into g occurred. All the a9-martensite peaks finally disappeared. It should be noted that the g-austenite peaks hardly shifted before the a9-martensite peaks disappeared. It should be considered that the decrease in diffraction intensity of a9-martensite might be related to sputter removal of the a9-martensitic layer on the surface. In order to evaluate the effect of sputtering, the mass loss during nitriding was measured with an analytical balance (accuracy 0.01 mg) on a nitrided area of 20=18 mm2. The data measured are shown in Table 1. Table 1 shows that there is actual mass loss, which means that sputter removal does occur during nitriding. If sputter removal of 1-mm-thick layer is assumed, the mass loss can be approximately calculated as follows:
Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
127
Fig. 2. XRD patterns of samples with an a9 layer nitrided for different times at 400–420 8C.
The area nitrided=the sputter depth=density of Fe s2.0 cm=1.8 cm=1=10y4 cm=7870 mgycm3 (2.833 mg. This suggests that the sputter depth during nitriding for 4 h could be approximately 0.12 mm according to the calculation above. If nitrogen incorporation is considered, the removal should be greater. All in all, sputter removal effects on the present transformation study can be neglected. The volume fractions of both phases transformed during the nitriding process can be calculated from quantitative XRD analysis using the following formula w8 x : V gs
1 Ia9 kg 1q Ig ka9
100% ;
Va9s1yVg
(1)
where Ia9 and Ig are the peak intensity of a9 and g, respectively, at any given moment; and ka9 and kg are
expressed as follows: w
z 1 2 )F) PwŽu.ey2M| ; 2 y v0g ~g
kgsx
w
z 1 2 )F) PwŽu.ey2M| 2 y v0a9 ~a9
ka9sx
and are dependent on u, hkl and the material measured, where v0g and v0a9 represent the unit cell volume of g and a9, respectively; P is the multiplicity factor of the hkl diffracting plane; w(u) is the Lorentz polarization (L-P) factor; F is the structure factor; and ey2M is the temperature factor. The XRD parameters above are listed in Table 2. It should be noted that the peak position of a9 and g have no obvious shift within a nitriding time of 0–90 min, although nitrogen atoms should be incorporated in lattice, which means that the lattice parameters of both phases do not obviously increase. The cell volumes of
Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
128
Table 2 XRD parameters of g and a9 phases (with CuKa) Phase
Parameter hkl
g a9
111 110
w(u) 11.8 11.3
v20 (nm) y2
4.6=10 2.4=10y2
±F±2 2
16 f 4f2
P
ey2Ma
k
8 6
0.979 0.979
6.9=105 f 2 4.7=105 f 2
f is the atomic scattering factor of Fe. w 9 x.
a
the a9 and g phases (v0a9, v0g) were calculated from the XRD results (ags0.3585 nm, aa9s0.2866 nm). According to the parameters in Table 2, kg yka9s1.47. The diffraction intensity measured and kg yka9s1.47 were substituted into Eq. (1), and the volume fractions of both phases transformed within the X-ray penetration of 2.5 mm can be obtained. The volume fractions of g and a9 transformed are plotted against nitriding time in Fig. 3. It is evident that the amount of a9 phase gradually decreased, but the g phase increased with increasing nitriding time. In order to clarify whether the transformation results from heating at the nitriding temperature or dissolution of nitrogen atoms in the a9 phase, the sample was vacuum-heated at 400–420 8C without the nitrogen atmosphere. This indicates that there is almost no change in the phase constituents and the diffraction intensity of both phases compared with Fig. 1b. This suggests that the transformation a9™g is not dependent on heating at the nitriding temperature. According to the vertical-section phase diagram of the Fe–Cr–Ni–C system at 8% Ni–18% Cr w10x, the a9 phase would transform to g phase only when the temperature reaches over 500 8C for 304 stainless steel (0.07 wt.% C). However, similar to carbon, nitrogen is an element that stabilizes the g phase, or enlarges the g-phase zone. The transformation temperature of a9™g would decrease with increasing nitrogen content. During plasma-nitriding, a lot of nitrogen atoms are dissolved in the a9 phase. When the nitrogen content in the a9 phase reaches a certain value to make the transformation temperature of a9™g lower than the nitriding temperature, the a9 phase transforms into the g phase. We conclude that the dissolution of nitrogen in a9 results in the transformation of a9™g, apart from the heating process at the nitriding temperature. According to the XRD results, the transformation of g™gN seems not to occur until the transformation of a9™g is complete. This means that nitrogen diffusion and dissolution preferentially occur in the a9 phase. It is well known that diffusion coefficient for nitrogen atoms in the a9 phase is 50-fold greater than in the closely stacked g phase. In addition, a9 induced by mechanical polishing should be mainly on the outermost surface. Nitrogen atoms diffuse in the a9 phase more easily than in the g phase. It is puzzling that the reverted
g phase on the outermost surface seems not transform to gN before the a9 phase disappears according to the XRD patterns. On the other hand, much more nitrogen can be in solid solution in the g lattice than in the a9 lattice. This suggests that when the nitrogen content in the a9 phase has become sufficient to induce the transformation of a9™g, nitrogen atoms dissolved in the g phase are not sufficient enough to observe the lattice expansion of gN. Moreover, when the transformation of a9™g occurs, two unit cells of a9 transform into one unit cell of g, which means that a volume contrast occurs. The transformation of a9™g is accompanied by a decrease in the free energy of the system. However, when the transformation of g™gN takes place, a volume expansion also occurs, so that the free energy of the system increases. An increase or decrease in the free energy of the system caused by both transformations may decide which transformation preferentially occurs. During completion of the transformation a9™g, nitrogen atoms are continuously dissolved and supersaturated in the g phase with treatment time. The transformation of g™gN would then occur, which can be identified by the peaks to the smaller-angle side of the g-austenite peaks (in Fig. 2, ts120 min). These peaks are associated with expanded austenite caused by the nitrogen remaining in solid solution in the f.c.c. lattice. The peaks were considerably broadened, which results from the inhomogeneous stress by supersaturating nitrogen and the non-uniform nitrogen content over the XRD depth range, as described in w11x. As the nitrogen content increases with treatment time, the expansion and broadening also increase. The plane spacing of gN(111) increased from 0.2107 to 0.2278 nm in a treatment time of 120–210 min. The SEM morphology of the nitrided sample layer for different times is shown in Fig. 4, with the thickness increasing over time. Although no lattice expansion in g was observed before the disappearance of the a9 peaks in our XRD experiments, the transformation g™gN may occur within a nitriding time of 90–120 min, because the expansion
Fig. 3. Volume fractions of a9 and g transformed during nitriding.
Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
129
2). Thus, the amount of lattice expansion for g in both is extremely different for the same nitriding time. The former (d111s0.2271 nm, ts120 min) is much greater than the latter (d111s0.2107 nm, ts120 min). The presence of an a9-martensite layer on the surface of austenitic stainless steel does not affect the last phase constituent, but the plasma-nitriding period needed is longer. The comparison experiments above also indirectly demonstrate that a two-step formation process for the plasma-nitrided layer described above is reasonable. 4. Conclusions
1. When plasma-nitriding austenitic stainless steel with an a9-martensitic layer induced by strain, the formation process for the nitride layer occurs in two steps. The a9-martensitic layer would preferably transform
Fig. 4. SEM morphology of samples nitrided for different times: (a) ts120; (b) ts15; and (c) ts210 min.
of g to gN appears to be a continuous process. However, the amount of lattice expansion should less than that observed for nitriding for 120 min, and might not be measured by XRD experiments. This should be further investigated. In any case, it can at least be concluded that when plasma-nitriding metastable austenitic stainless steel with an a9-martensitic layer, the process of formation of the nitrided layer takes two steps, a9™g transformation, followed by the reverted g™gN transformation. XRD patterns of the electric-polished sample nitrided for different times are shown in Fig. 5. It is evident that gN peaks appear after nitriding for 30 min, whereas gN peaks do not appear until after 120 min when nitriding a sample with a pre-existing a9-martensite layer (Fig.
Fig. 5. XRD patterns of samples without an a9 layer nitrided for different times at 400–420 8C.
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Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
to g, followed by the transformation g™gN. The volume fractions of a9 and g transformed within the X-ray penetration depth of 2.5 mm have been calculated by quantitative XRD analysis. 2. The amount of lattice expansion and the thickness of the nitrided layer (gN) increase with time. However, for the same nitriding time, the amount of the lattice expansion of gN for the sample with a pre-existing a9-martensite layer is much less than that of the sample without an a9 layer, which indirectly demonstrates that the two-step transformation process actually occurs. 3. Transformation of a9™g is dependent on the nitrogen content in a9, other than resulting from heating at the plasma-nitriding temperature. 4. Although gN peaks did not appear until the a9 peaks had disappeared in the present study, the transformation of the reverted g™gN may occur within the nitriding time of 90–120 min before complete transformation of a9™g.
References w1x X.L. Xu, L. Wang, Z.W. Yu, Z.K. Hei, Metall. Trans. A 27 (1996) 1347. w2x X.L. Xu, L. Wang, Z.W. Yu, Z.K. Hei, Surf. Coat. Technol. 132 (2–3) (2000) 270. w3x C. Blawert, A. Weisheit, B.L. Mordike, F.M. Knoop, Surf. Coat. Technol. 85 (1996) 15. w4x Z.L. Zhang, T. Bell, Surf. Eng. 1 (1985) 131. w5x L. Wang, X.L. Xu, Z.W. Yu, Z.K. Hei, Surf. Coat. Technol. 124 (2000) 93. w6x L. Wang, B. Xu, Z. Yu, Y. Shi, Surf. Coat. Technol. 130 (2000) 304. w7x R. Wei, J.J. Vajo, J.N. Matossian, P.J. Wilbur, J.A. Daris, D.L. Williamson, Surf. Coat. Technol. 83 (1996) 235. w8x B.D. Cullity, Elements of X-Ray Diffraction, Addison-Wesley, USA, 1978, p. 412. w9x A Buerger, et al., International Tables for X-Ray Crystallography, III, Kynoch press, Birmingham, England, 1976, p. 232. w10x E.E. Thum, The Book of Stainless Steels, ASM, 1933, p. 42. w11x O. Ozturk, D.L. Williamson, J. Appl. Phys. 77 (1995) 3839.
Structural characteristics of low-temperature plasma-nitrided layers on AISI 304 stainless steel with an a9-martensite layer Zhiwei Yu, Xiaolei Xu*, Liang Wang, Jianbing Qiang, Zukun Hei Dalian Maritime University, Institute of Metal and Technology, Dalian 116024, PR China Received 7 March 2001; accepted in revised form 1 November 2001
Abstract The characteristics of the structure formation of a low-temperature plasma-nitrided layer on AISI 304 austenitic stainless steel with a pre-existing a9-martensitic deformation layer were studied in detail by glancing-angle X-ray diffraction (XRD). The results show a two-step formation process for the plasma-nitrided layer. The a9-martensitic layer gradually reverts to the original gaustenite structure at first, followed by the transformation of g™gN as the nitrogen content increases with treatment time. Within the same nitriding time, the sample with an a9 layer shows much less lattice expansion than the sample without an a9 layer. The volume fractions of a9 and g phases transformed during the plasma-nitriding process were calculated by quantitative XRD. The transformation of a9™g is dependent on the nitrogen content dissolved in a9, apart from that resulting from heating at the nitriding temperature. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Plasma nitriding; Stainless steel; Transformation
1. Introduction Most stainless steels (and in particular austenitic grades) have relatively low hardness, and therefore poor resistance to wear, unless a surface modification treatment can be applied. Plasma nitriding has been widely used in industrial production to increase the surface hardness, and hence improve the wear resistance of ferrous materials. Generally, the treatment temperature for traditional plasma-nitriding is above 500 8C w1x in order to obtain a sufficiently thick modified layer. However, CrN precipitation w2,3x occurs for most stainless steels when the nitriding temperature is increased above 450 8C, so that their corrosion resistance is largely decreased. Reducing the treatment temperature can avoid precipitation of CrN and a nitrogen-expanded austenite (gN) can form, but lower nitriding efficiency results, leading to increased treatment times w4x. Traditional plasma-nitriding technologies are therefore of limited use in industrial treatment of stainless steels. A new, low-pressure plasma arc-source ion nitriding technology developed by the authors w5,6x can substantially increase *Corresponding author. Tel.: q86-411-4729613; fax: q86-4114671395. E-mail address: [email protected] (X. Xu).
the efficiency of plasma-nitriding at low temperature and an adequately thick gN modified layer can be obtained in a shorter treatment time. There are many reports on gN-modified layers on austenitic stainless steels produced by plasma-nitriding, and plasma immersion (PIII) and plasma source ion implantation (PSII) w3,7x, but in most experiments, an ‘ideal’ specimen surface having a single g-phase structure was selected, which does not correspond to the typical production conditions. Usually, there is a hardened layer formed on the surface of metastable austenitic stainless steel by machining, which consists of the deformation martensite (a9) coexisting with g-austenite. There should be a difference in the structural characteristics and formation process for the nitrided layer for austenitic stainless steel with or without a pre-introduced a9 layer. In this paper, we investigate the structural characteristics of nitrided layers on 304 stainless steels with an a9 deformation layer and examine the related transformation during the plasma-nitriding process. 2. Experimental methods The composition of the annealed AISI 304 austenite stainless steel used for this work was (in wt.%): C,
0257-8972/02/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 1 . 0 1 6 7 0 - X
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Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
0.07; Si, 0.78; Mn, 0.90; S, 0.13; P, 0.03; Cr, 19.0; Ni, 8.26; and Fe, balance. Plate samples of dimensions 15=15=5 mm3 were prepared and mechanically polished using 800噛 grit SiC paper. The samples were plasma-nitrided using the above-mentioned low-pressure plasma arc source at 400–420 8C. Ammonia was used as the working gas. The working pressure was approximately 0.4 Pa. The arc current was 40 A with a voltage of 50 V. A negative bias of 0.8–1 kV was applied to the substrates. The current density on the substrate surface was 0.6–0.8 mAycm2 The phases present in the modified layer were determined by glancing-angle Xray diffraction on a Rigaku Dymax-IIIA diffractometer. The radiation used was CuKa (ls0.1542 nm) at an incident angle of 108. The X-ray penetration depth under these diffraction conditions was estimated to be approximately 2.5 mm. 3. Results and discussion 3.1. Structure of the matrix XRD patterns of the electrically and mechanically polished samples are shown in Fig. 1. The XRD pattern in Fig. 1a is typical of a f.c.c. structure, the peaks of which appear at 2us43.718, 50.928, 74.838, etc., corresponding to plane spacings of ds0.2069, 0.1792, 0.128
Table 1 Mass change of the sample during nitriding Nitriding
Mass (g)
time
Measured
Average
6.61909 6.61909 6.61909 6.61909 6.61907 6.61906 6.61907 6.61906 6.61904 6.61905 6.61904 6.61905 6.61874 6.61875 6.61875 6.61874
6.61909
0
12 min
18 min
4h
Mass loss (mg)
6.61907
0.02
6.61905
0.04
6.61875
0.34
nm, etc., respectively. This demonstrates that the electrically polished sample is composed of single-phase gaustenite. However, in the XRD pattern of Fig. 1b, the peaks (2us44.568, 64.728, 82.728) of the b.c.c. structure corresponding to the plane spacings of the ferrite (ds 0.2032, 0.1440, 0.1171 nm) appear, except for the peaks of g-austenite. Its peak intensity is higher than that of g-austenite, which demonstrates that there is a certain amount of a9-martensite induced by the strain on the surface during mechanical polishing of the sample. The peaks of the remaining g-austenite did not shift. 3.2. Structure composition of the plasma-nitrided layer
Fig. 1. XRD patterns of 304 stainless steel (a) electrically and (b) mechanically polished.
XRD patterns of samples nitrided for different times are shown in Fig. 2. It is evident that the peak intensity of a9-martensite gradually decreased, but the intensity of the g-austenite increased with treatment time. This suggests that a gradual transformation of a9 into g occurred. All the a9-martensite peaks finally disappeared. It should be noted that the g-austenite peaks hardly shifted before the a9-martensite peaks disappeared. It should be considered that the decrease in diffraction intensity of a9-martensite might be related to sputter removal of the a9-martensitic layer on the surface. In order to evaluate the effect of sputtering, the mass loss during nitriding was measured with an analytical balance (accuracy 0.01 mg) on a nitrided area of 20=18 mm2. The data measured are shown in Table 1. Table 1 shows that there is actual mass loss, which means that sputter removal does occur during nitriding. If sputter removal of 1-mm-thick layer is assumed, the mass loss can be approximately calculated as follows:
Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
127
Fig. 2. XRD patterns of samples with an a9 layer nitrided for different times at 400–420 8C.
The area nitrided=the sputter depth=density of Fe s2.0 cm=1.8 cm=1=10y4 cm=7870 mgycm3 (2.833 mg. This suggests that the sputter depth during nitriding for 4 h could be approximately 0.12 mm according to the calculation above. If nitrogen incorporation is considered, the removal should be greater. All in all, sputter removal effects on the present transformation study can be neglected. The volume fractions of both phases transformed during the nitriding process can be calculated from quantitative XRD analysis using the following formula w8 x : V gs
1 Ia9 kg 1q Ig ka9
100% ;
Va9s1yVg
(1)
where Ia9 and Ig are the peak intensity of a9 and g, respectively, at any given moment; and ka9 and kg are
expressed as follows: w
z 1 2 )F) PwŽu.ey2M| ; 2 y v0g ~g
kgsx
w
z 1 2 )F) PwŽu.ey2M| 2 y v0a9 ~a9
ka9sx
and are dependent on u, hkl and the material measured, where v0g and v0a9 represent the unit cell volume of g and a9, respectively; P is the multiplicity factor of the hkl diffracting plane; w(u) is the Lorentz polarization (L-P) factor; F is the structure factor; and ey2M is the temperature factor. The XRD parameters above are listed in Table 2. It should be noted that the peak position of a9 and g have no obvious shift within a nitriding time of 0–90 min, although nitrogen atoms should be incorporated in lattice, which means that the lattice parameters of both phases do not obviously increase. The cell volumes of
Z. Yu et al. / Surface and Coatings Technology 153 (2002) 125–130
128
Table 2 XRD parameters of g and a9 phases (with CuKa) Phase
Parameter hkl
g a9
111 110
w(u) 11.8 11.3
v20 (nm) y2
4.6=10 2.4=10y2
±F±2 2
16 f 4f2
P
ey2Ma
k
8 6
0.979 0.979
6.9=105 f 2 4.7=105 f 2
f is the atomic scattering factor of Fe. w 9 x.
a
the a9 and g phases (v0a9, v0g) were calculated from the XRD results (ags0.3585 nm, aa9s0.2866 nm). According to the parameters in Table 2, kg yka9s1.47. The diffraction intensity measured and kg yka9s1.47 were substituted into Eq. (1), and the volume fractions of both phases transformed within the X-ray penetration of 2.5 mm can be obtained. The volume fractions of g and a9 transformed are plotted against nitriding time in Fig. 3. It is evident that the amount of a9 phase gradually decreased, but the g phase increased with increasing nitriding time. In order to clarify whether the transformation results from heating at the nitriding temperature or dissolution of nitrogen atoms in the a9 phase, the sample was vacuum-heated at 400–420 8C without the nitrogen atmosphere. This indicates that there is almost no change in the phase constituents and the diffraction intensity of both phases compared with Fig. 1b. This suggests that the transformation a9™g is not dependent on heating at the nitriding temperature. According to the vertical-section phase diagram of the Fe–Cr–Ni–C system at 8% Ni–18% Cr w10x, the a9 phase would transform to g phase only when the temperature reaches over 500 8C for 304 stainless steel (0.07 wt.% C). However, similar to carbon, nitrogen is an element that stabilizes the g phase, or enlarges the g-phase zone. The transformation temperature of a9™g would decrease with increasing nitrogen content. During plasma-nitriding, a lot of nitrogen atoms are dissolved in the a9 phase. When the nitrogen content in the a9 phase reaches a certain value to make the transformation temperature of a9™g lower than the nitriding temperature, the a9 phase transforms into the g phase. We conclude that the dissolution of nitrogen in a9 results in the transformation of a9™g, apart from the heating process at the nitriding temperature. According to the XRD results, the transformation of g™gN seems not to occur until the transformation of a9™g is complete. This means that nitrogen diffusion and dissolution preferentially occur in the a9 phase. It is well known that diffusion coefficient for nitrogen atoms in the a9 phase is 50-fold greater than in the closely stacked g phase. In addition, a9 induced by mechanical polishing should be mainly on the outermost surface. Nitrogen atoms diffuse in the a9 phase more easily than in the g phase. It is puzzling that the reverted
g phase on the outermost surface seems not transform to gN before the a9 phase disappears according to the XRD patterns. On the other hand, much more nitrogen can be in solid solution in the g lattice than in the a9 lattice. This suggests that when the nitrogen content in the a9 phase has become sufficient to induce the transformation of a9™g, nitrogen atoms dissolved in the g phase are not sufficient enough to observe the lattice expansion of gN. Moreover, when the transformation of a9™g occurs, two unit cells of a9 transform into one unit cell of g, which means that a volume contrast occurs. The transformation of a9™g is accompanied by a decrease in the free energy of the system. However, when the transformation of g™gN takes place, a volume expansion also occurs, so that the free energy of the system increases. An increase or decrease in the free energy of the system caused by both transformations may decide which transformation preferentially occurs. During completion of the transformation a9™g, nitrogen atoms are continuously dissolved and supersaturated in the g phase with treatment time. The transformation of g™gN would then occur, which can be identified by the peaks to the smaller-angle side of the g-austenite peaks (in Fig. 2, ts120 min). These peaks are associated with expanded austenite caused by the nitrogen remaining in solid solution in the f.c.c. lattice. The peaks were considerably broadened, which results from the inhomogeneous stress by supersaturating nitrogen and the non-uniform nitrogen content over the XRD depth range, as described in w11x. As the nitrogen content increases with treatment time, the expansion and broadening also increase. The plane spacing of gN(111) increased from 0.2107 to 0.2278 nm in a treatment time of 120–210 min. The SEM morphology of the nitrided sample layer for different times is shown in Fig. 4, with the thickness increasing over time. Although no lattice expansion in g was observed before the disappearance of the a9 peaks in our XRD experiments, the transformation g™gN may occur within a nitriding time of 90–120 min, because the expansion
Fig. 3. Volume fractions of a9 and g transformed during nitriding.
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2). Thus, the amount of lattice expansion for g in both is extremely different for the same nitriding time. The former (d111s0.2271 nm, ts120 min) is much greater than the latter (d111s0.2107 nm, ts120 min). The presence of an a9-martensite layer on the surface of austenitic stainless steel does not affect the last phase constituent, but the plasma-nitriding period needed is longer. The comparison experiments above also indirectly demonstrate that a two-step formation process for the plasma-nitrided layer described above is reasonable. 4. Conclusions
1. When plasma-nitriding austenitic stainless steel with an a9-martensitic layer induced by strain, the formation process for the nitride layer occurs in two steps. The a9-martensitic layer would preferably transform
Fig. 4. SEM morphology of samples nitrided for different times: (a) ts120; (b) ts15; and (c) ts210 min.
of g to gN appears to be a continuous process. However, the amount of lattice expansion should less than that observed for nitriding for 120 min, and might not be measured by XRD experiments. This should be further investigated. In any case, it can at least be concluded that when plasma-nitriding metastable austenitic stainless steel with an a9-martensitic layer, the process of formation of the nitrided layer takes two steps, a9™g transformation, followed by the reverted g™gN transformation. XRD patterns of the electric-polished sample nitrided for different times are shown in Fig. 5. It is evident that gN peaks appear after nitriding for 30 min, whereas gN peaks do not appear until after 120 min when nitriding a sample with a pre-existing a9-martensite layer (Fig.
Fig. 5. XRD patterns of samples without an a9 layer nitrided for different times at 400–420 8C.
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to g, followed by the transformation g™gN. The volume fractions of a9 and g transformed within the X-ray penetration depth of 2.5 mm have been calculated by quantitative XRD analysis. 2. The amount of lattice expansion and the thickness of the nitrided layer (gN) increase with time. However, for the same nitriding time, the amount of the lattice expansion of gN for the sample with a pre-existing a9-martensite layer is much less than that of the sample without an a9 layer, which indirectly demonstrates that the two-step transformation process actually occurs. 3. Transformation of a9™g is dependent on the nitrogen content in a9, other than resulting from heating at the plasma-nitriding temperature. 4. Although gN peaks did not appear until the a9 peaks had disappeared in the present study, the transformation of the reverted g™gN may occur within the nitriding time of 90–120 min before complete transformation of a9™g.
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