Formation of chromium nitride coatings on carbon steels by pack cementation process

SCT-21692; No of Pages 7 Surface & Coatings Technology xxx (2016) xxx–xxx

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Formation of chromium nitride coatings on carbon steels by pack cementation process X.J. Lu, Z.D. Xiang ⁎ The State Key Laboratory of Refractories and Metallurgy, School of Materials and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, PR China

a r t i c l e

i n f o

Article history: Received 9 August 2016 Revised 16 October 2016 Accepted in revised form 17 October 2016 Available online xxxx Keyword: Nitriding Chromising Carbon steel Pack cementation Chromium nitride

a b s t r a c t Chromium nitride coatings are usually applied to metal surface by physical vapour deposition processes, which not only have line-of-sight restrictions but also require capital intensive facilities. This study is thus undertaken in an attempt to form chromium nitride coatings on the surface of carbon steels by pack cementation process, which offers an efficient method to form chromium nitride coatings on metal surfaces of complex geometries. Thermal chemical calculations were undertaken to identify suitable activators for the process. Experiments were then carried out using pack powder mixture of a composition 30Cr2N-2NH4Cl-68Al2O3 (wt.%). Two grades of carbon steels with compositions differing mainly in carbon content were used as substrates. The process was carried out in the temperature range of 1000 °C–1100 °C. The microstructures of the coatings were characterised using scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS). The phases in the coatings were analysed by X-ray diffraction (XRD). In all cases, the coatings produced had an outer Cr2N layer, but the inner microstructures depended on carbon content in steel substrate. The effects of processing temperature and time on the microstructures of the coatings were also discussed. © 2016 Elsevier B.V. All rights reserved.

1. Introduction Chromium nitride coatings in the form of either CrN or Cr2N or a combination of both have been applied to surfaces of different grades of steels, usually to increase their surface hardness and hence tribological performance [1–7]. More recently, they have also been investigated for use as corrosion resistant coatings on metal alloy bipolar plates in proton exchange membrane fuel cells where the coating must also have low interfacial electrical contact resistance [8–11]. The processes used most often for depositing this type of coatings are physical vapour deposition processes such as magnetron sputtering [1–5] and cathodic arc deposition [6–7]. High quality coatings free of defects such as voids can be deposited using these methods. But these techniques require capital intensive facilities and, in addition, they all have line-of-sight restrictions. On the other hand, nitriding is a thermochemical surface treatment technique that has no such limitations and as such it has also been investigated for forming chromium nitride coatings on metal surfaces. But, in the case of steels, nitriding process cannot be used alone because the Cr concentration in nearly all commercial steels, including various grades of high Cr stainless steels, is not high enough to enable the formation of an external chromium nitride layer at high nitriding temperatures. To obtain an external chromium nitride layer on steel surfaces, nitriding process has to be used in combination with ⁎ Corresponding author. E-mail address: [email protected] (Z.D. Xiang).

other processes such as Cr-plating or chromising, i.e., the Cr concentration in the steels' surface has to be raised first by the application of one of these processes before nitriding process can be applied [12–16]. Only in the case of special alloys that contain sufficiently high Cr concentration such as Ni-50Cr, can an external chromium nitride surface layer be formed via selective nitridation [8,17–18]. But, even for Ni-50Cr alloy, only a thin chromium nitride layer can be obtained with a nonuniform thickness varying from a minimum of approximately 1 μm to a maximum of roughly 5 μm, although such a thin chromium nitride layer was found to be sufficient to provide the required corrosion resistance in corrosive environments in proton exchange membrane fuel cells [19]. Although the process of combining nitriding with prior Cr-plating or chromising as described above is technically feasible for forming a chromium nitride coating on steel surface, such a two step process may be suitable only for applications where cost is of less concern; a single step process would be highly desirable. Pack cementation could offer such a possibility as it is a process by which more than one element may be simultaneously diffused into metal surfaces to form diffusion coatings [20–23]. In a previous study, the feasibility of nitriding the surface of austenitic stainless steels in un-activated packs was demonstrated [24], and in a subsequent study, it was further demonstrated that, with activated packs, the process can be applied to form a Cr2N top layer coating with a Cr-enriched layer underneath on the surface of austenitic stainless steel [25]. In the present study, the viability of the same process for forming chromium nitride coatings on carbon steels will be

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0

investigated, and the effects of processing conditions and carbon content in steel substrates on the microstructure of the coatings produced will be delineated.

(a)

Cr2 N ¼ Cr þ N2

ð1Þ

Cr produced in the above reaction will in turn react with chloride in the pack to generate various Cr-chloride vapour species, which would include CrCl, CrCl2, CrCl3 and Cr2Cl6. Among these species of Cr-chloride vapours, only CrCl functions as the species that deposits Cr on the metal surface; other types of Cr-chloride vapour species merely participate in the vapour transportation processes from the pack powder to metal surface [29–30]. Therefore, the partial pressure of CrCl vapour species may be used to compare the activating strength of different chloride salts. A higher CrCl partial pressure would suggest a higher activating strength

-4 NH4Cl pack AlCl3 pack

-6

NaCl pack -8 -3

(b) -6 log p CrCl (atm)

For the simple purpose of pack nitriding, various types of nitride powers may be used to serve as a source of nitrogen in powder packs. In this case, the packs may not need to be activated as the nitrogen partial pressures from partial dissociation of nitride powders can be sufficiently high for nitrogen enrichment in metal surfaces. These nitrogen partial pressures can be readily calculated using standard thermodynamic data or commercially available software for thermochemical calculations for commonly available nitride powders such as Cr2N, Si3N4 and TiN. The calculations have shown that the Cr2N power is most suitable for pack nitriding as it generates the highest nitrogen partial pressure among the types of nitride powders listed above [24]. Thus, before commencing experiments, it was considered that the Cr2N powder may also be used to form a chromium nitride coating on metal surfaces in pack cementation process if the pack used for the process is activated by a suitable activator, which can be a chloride or fluoride salt. The condition necessary for forming such a coating is that the packs must be able to generate sufficiently high partial pressures of both nitrogen and Cr-chloride or Cr-fluoride vapours at processing temperatures. The selection of halide activators to fulfill this condition can be made on the basis of thermochemical calculations, which may be carried out with the assistance of a thermochemical calculation software. In this study, only chloride salts are considered for use as potential activators for forming a chromium nitride coating in packs containing Cr2N powder. The commonly used chloride salts for pack cementation process are NH4Cl, NaCl and AlCl3, which are the salts considered here. For an efficient pack cementation process, the amount of halide activator added is normally controlled in the range of 1 to 5 wt.%. Thus, the amount of halide salts addition in the packs was selected at 2 wt.%. To achieve high coating formation rate, the amount of powder for the source of coating formation elements should be as high as possible, but not too high to prevent powder sintering at high temperatures [26–28]. For this reason, the amount of Cr2N powder selected was 30 wt.%. Thus, thermochemical calculations were carried out for pack powder mixtures of composition 30Cr2N-2Halide salt-68Al2O3 (wt.%). The calculations were made using ChemSage software in combination with the pure substance data base SPS on the condition that total pressure in the packs is 1 atm. The calculation results showed that the N2 partial pressure in NH4Cl activated pack is the highest among the packs considered here (Fig. 1a). It can be noted that the N2 partial pressures in AlCl3 and NaCl activated packs are the same at all the temperatures. In fact, further analysis showed that the N2 partial pressures in the latter two packs are also the same as in un-activated Cr2N powder pack, indicating that AlCl3 and NaCl do not react with Cr2N powder at elevated temperatures and the N2 partial pressures generated in AlCl3 and NaCl activated packs are simply the results of Cr2N powder dissociation according to

log p N2 (atm)

-2 2. Basis of activator selection for pack chromium nitride coating formation

-9 NH4Cl pack AlCl3 pack

-12

NaCl pack -15 600

700

800

900

1000

1100

1200

Temperature (˚C) Fig. 1. Partial pressures of N2 and CrCl calculated for packs 30Cr2N-2Halide-68Al2O3 (wt.%), where Halide is NH4Cl or AlCl3 or NaCl; (a) N2, (b) CrCl.

as it would lead to a higher rate of Cr deposition and consequently result in faster Cr enriched surface layer growth. Thus, the comparison of the calculated partial pressures of CrCl vapour species as shown in Fig. 1b suggests that, in terms of chromising process, the activating strength of AlCl3 is slightly higher than that of NH4Cl, and the activating strengths of these two types of chloride salts are all much higher than that of NaCl. Thus, for chromising purpose, NH4Cl would be a slightly weaker activator than AlCl3. But, the N2 partial pressure generated in NH4Cl activated packs is much higher than in AlCl3 activated packs, which means that the conditions for chromium nitride coating formation in the former packs are much more favourable than in the latter packs. Therefore, NH4Cl was chosen as the activator for coating experiments in this study. 3. Materials and experimental details Two grades of carbon steels, 45# and Q235C, were used as substrate materials in this study. The steel 45# is a grade with a composition range of 0.50–0.80Mn, Cr ≤ 0.25, Ni ≤ 0.30, Cu ≤ 0.25, 0.17–0.37Si and 0.42–0.5C (wt.%), and for steel Q235C, the composition range is specified as 0.35–0.80Mn, Si ≤ 0.35 and C ≤ 0.17 (wt.%). The reason for including the latter steel grade in this study was to investigate the effects of carbon content in steel substrates on the microstructures of the coatings formed on the steel surfaces. Steel plates were cut and ground to an approximate dimension of 20 × 10 × 2 mm. In all cases, the pack powder composition used for coating experiments was 30Cr2N–2NH4Cl– 68Al2O3 (wt.%), which is the same as that used in thermochemical calculations as described in the previous section. The particle sizes of Cr2N (commercial grade, Jinzhou Metal Research Institute, China) and Al2O3 powders (Acros, 99%, CAS:1344-28-1) were all below 40 μm. NH4Cl was manually ground in a ceramic mortal with a ceramic pestle; the powder obtained was not sieved and used as it was.

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The specimens were cleaned with organic solvent and then buried in the pack powder in a cylindrical alumina crucible of 30 mm in diameter and 40 mm in length. The specimen was buried in a longitudinal direction in the crucible, i.e., with its length in the same direction as that of the longitudinal axis of the cylindrical crucible. Only one specimen was packed into the crucible and care was taken to ensure that the location of the specimen in the crucible was approximately the same in each experiment. The pack powder was filled right to the top of crucible and the pack was vibrated several times during filling process to consolidate the pack powder. The crucible was then covered with an alumina lid and sealed from the outside using a high temperature cement. It is useful to mention that the seal applied between crucible and lid is not gas-tight. No attempt was made to control the humidity in the packs. The coating process was carried out in a tube furnace in the temperature range of 1000–1100 °C. Flowing Ar was maintained during heating, dwelling and cooling processes to ensure a protective atmosphere throughout the coating process. The heating rate used was 10 °C/min and all the specimens were furnace-cooled after dwelling at the coating temperature for a specified length of time. The coating times reported here were the dwelling times at the coating temperatures. Microstructures of coated specimens were analysed using field emission scanning electron microscopy (SEM) (FEI Quanta 200) and energy dispersive spectroscopy (EDS) (Oxford INCA). The phases formed in the surface layer of treated specimens were analysed by X-ray diffraction (XRD) (PHILIPS X'Pert PRO MPD diffractometer with Cu-Kα source radiation operated at 40 kV and 40 mA). 4. Results and discussion

3

small amount of Fe, and a small amount of carbon may also be present in it. It is also evident that the Cr2N coating layer resulted from a highly oriented growth with the (002) reflection showing the highest intensity. For a non-textured Cr2N coating layer, the highest peak intensity is expected to be from the (111) reflection as shown by the XRD pattern measured for Cr2N powder (Fig. 2), but, it is one of the weakest among all the reflection peaks in the XRD pattern measured from the Cr2N coating layer. This oriented growth feature for the Cr2N coating layer produced in the pack process differs from that of the Cr2N coating formed in magnetron sputtering process in which it grew predominantly to have a (111) orientation [3]. Whether this difference in texture is mainly due to the difference in the steel grades used as substrates by different investigators is not clear at present and requires further investigation. The cross-sectional SEM image and the element depth profiles measured by EDS showed that the coating had a multiple layer structure (Fig. 3). The outmost layer was Cr2N as discussed above. This layer was quite uniform in thickness, which was approximately 17 μm. A small amount of Fe was present in this layer (Fig. 3b), which was likely to be the cause for shifting the X-ray reflection peaks of the Cr2N layer to the right (higher angles) (Fig. 2). Nevertheless, the residual stresses resulting from cooling due to thermal expansion difference are likely to be present in the Cr2N layer, which may also contribute to the XRD peak shift observed. Underneath the outmost Cr2N layer, there was a visually identifiable layer with a thickness of about 6 μm. This inner layer contained Cr and Fe but no nitrogen (Fig. 3b). The Cr concentrations across this layer were all above 75 at.%, and the Fe concentrations were all below 25 at.% (Fig. 3b). In order to identify the phase of this layer, the outmost

4.1. Phase and microstructure of the coating

(a)

Experiments were carried out initially at 1100 °C using steel 45# as substrate. The specimen obtained after coating at this temperature for 4 h had a quite smooth surface with a bright and shining metallic appearance, and the specimen weight gain was 11.4 mg/cm2, indicating that a surface layer was produced. No entrapped pack powder particles were present in the as-coated surface. A simple comparison of XRD pattern measured directly from the ascoated surface with that from Cr2N powder confirmed that the phase of the outmost surface layer was Cr2N (Fig. 2). But, it can be noted that its reflection peaks were shifted slightly to the right relative to those of Cr2N powder, indicating that the chemical composition of the Cr2N coating layer differed slightly from that of the pure Cr2N. The EDS results to be presented later will show that the Cr2N coating layer contained a

(Cr, Fe)7C3

(b)

3000 2500

100

(002)

(113)

Concentration (at.%)

Coating

2000 Intensity

Cr2N

1500 (112)

(110) (111)

1000

(302)

(111)

500

Cr2N powder (221) (112) (300) (113) (302)

(002) (110)

0 30

40

50

60

70

80

90

2θ (˚)

80 Cr N

60 Cr-enriched interdiffusion zone

40

Fe

20 0 0

20

40

60

Depth (µm) Fig. 2. Comparison of XRD patterns measured from the as-coated surface with that measured from Cr2N powder ; the coating was formed on steel 45# by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at 1100 °C for 4 h. (PDF#: 00035-0803).

Fig. 3. Cross-sectional SEM and element depth profiles measured by EDS for a coating formed on steel 45# by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at 1100 °C for 4 h.

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Cr2N layer of a specimen coated also at 1100 °C but for only 2 h was manually ground away and, during the grinding process, microscopic checks were made to ensure that the outmost Cr2N layer was completely removed. The XRD pattern was then measured from the new surface. The result showed that this layer was essentially a chromium carbide, and in this case, it was of the type Cr7C3 (Fig. 4). Because this layer also contained Fe, all the reflection peaks were shifted slightly to higher angles as compared with the standard Cr7C3 reflection peaks in the XRD database. Thus, in effect, the phase of the inner layer was (Cr, Fe)7C3. It is apparent that the carbon concentration in the steel and Cr concentration deposited into the steel surface from the vapour phase in the pack at 1100 °C were high enough to have enabled the formation of this continuous inner layer of Cr-carbide below the outmost Cr2N layer. There was also a large interdiffusion zone between the inner Cr-carbide layer and steel substrate, the thickness of which was about 23 μm. This interdiffusion zone is not visible on the SEM micrograph but can be easily recognised on the Cr and Fe concentration depth profiles measured by EDS (Fig. 3b). This zone was enriched with Cr, but with its concentration showing a diffusion type of depth profile, decreasing progressively from approximately 24 at.% at a depth just below the (Cr, Fe)7C3 layer to the same level as that in the steel substrate as depth increases. This feature is consistent with that of Cr concentration depth profile normally observed for coatings formed in pack chromising process [31–32]. This is relevant because it indicates that Cr atoms diffused also across the Cr-carbide layer to have enabled the growth of this interdiffusion zone. The results above suggest that chromising and nitriding processes took place in situ in the pack, which led to the formation of a coating with a three-layer structure consisting of an outer Cr2N layer, an inner Cr-carbide inner layer and a Cr-enriched interdiffusion zone between the inner Cr-carbide layer and steel substrate. According to this coating structure, it is tempting to hypothesize that chromising, which involves decomposition of Cr-chloride vapours to deposit Cr on steel surface, i.e. 2CrCl = 2Cr + Cl2, took place first in the pack, which initially led to the formation of a Cr-enriched surface through inward Cr diffusion. Cl2 released is reacted again with Cr released in the pack through Eq. (1) to form Cr-chloride vapour species. When the Cr concentration in steel surface was increased to a high enough level, the reaction N2 + 4Cr (in solid solution) = 2Cr2N occurred in-situ, which resulted in the formation of a continuous Cr2N layer on the outmost surface. To verify whether this was the case and to assess the growth kinetics of both outer Cr2N and inner Cr-carbide layers, the effect of coating time on coating microstructure was investigated.

2000

(421)

Cr7C3

1500 Intensity

(202) 1000

(402)

500

(411)

(403)

(440)

(552) (741) (1001)

(801) 0 30

40

50

60

70

80

90

2θ Fig. 4. XRD pattern measured from a surface with only the outmost Cr2N layer being removed by grinding for a coating formed on steel 45# by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at 1100 °C for 2 h. (PDF#: 00-036-1482).

4.2. Dependence of coating microstructure on coating time The steel 45# was also used as the substrate in this section. The pack process was carried out at 1100 °C for all the specimens, but the coating time was varied from 0 h to 8 h. The cross sectional images for specimens with coating times varying from 0 h to 4 h are compared in Fig. 5. It shows that all the coatings obtained had the same cross-sectional microstructure, consisting of an outer layer and an inner layer. There was also a Cr-enriched diffusion zone underneath the inner layer in all cases, which, however, was not visually visible on the SEM images as discussed in the previous section. The XRD and EDS measurements suggested that the outer layer for all the specimens was Cr2N and the inner layer was Cr-carbide. But, the type of Cr-carbides varied with coating time. For instance, as shown in the previous section, the Cr-carbide in the inner layer was (Cr, Fe)7C3 for the specimen coated for 2 h, but, the same analysis showed (the result will be shown later) that it was (Cr, Fe)23C6 for the specimen coated for 8 h. Therefore, it appears that the type of Cr-carbides in the inner layer changed from a lower Cr concentration type to a higher Cr concentration type as coating time increased. The fact that both outer Cr2N layer and inner Cr-carbide layer were present in the specimen surface at 0 h coating time suggests that these two layers started to form simultaneously during heating stage and grew in thickness as coating time increased (Fig. 5). Detailed analysis showed that the total thickness of these two layers dT, i.e., dT = Cr2N layer thickness + Cr-carbide layer thickness, increased linearly with square root of coating time t or t1/2, although the slope of this linear relationship changed at approximately t = 2.7 h (Fig. 6a), suggesting that the growth of this total layer thickness was controlled by a diffusion mechanism. This is consistent with the coating growth mechanism observed in general for coatings formed in pack cementation process [26–28]. However, the growth behaviour of individual layers differed as shown in Fig. 6b, which plots the layer thickness versus coating time t for individual Cr2N and Cr-carbide layers. The growth of the outer Cr2N layer continued until coating time reached 4 h, after which the growth of this layer stopped. For the inner Cr-carbide layer, its thickness decreased abruptly at a coating time between 2 h and 4 h. Both below and above this coating time, the inner Cr-carbide layer thickness increased continuously with coating time. It can be seen that the inner Cr-carbide layer continued to grow at coating times longer than 4 h, which were the times when the outer Cr2N layer had stopped growing. Therefore, it seems that Cr atoms deposited from the vapour phase in the pack can diffuse across the outer Cr2N layer to enable the continued growth of the inner Cr-carbide layer even when the growth of the outer Cr2N layer has stopped. It is of importance to take notice of another feature of individual layer growth in the coatings, i.e., at coating times shorter than 2 h, the thickness of the outer Cr2N layer was smaller than that of the inner Cr-carbide layer, but at coating times longer than 4 h, the order was reversed (Fig. 6b). This change in layer thickness order happened at a coating time between 2 h and 4 h, which coincided with the coating time interval in which the inner Cr-carbide layer thickness showed a sudden decrease as described previously. It is thus highly likely that both of these two phenomena occurred at the same coating time. These phenomena may be accounted for by the same process, i.e., the change in the phase of the inner Cr-carbide layer as coating time increased. As discussed previously, the phase of inner Cr-carbide layer was (Cr, Fe)7C3 in the specimen coated for 2 h, but it changed to (Cr, Fe)23C6 in the specimen coated for 8 h as shown in Fig. 7, which was measured from a surface after manually grinding away the outer Cr2N layer (however, a strong peak at 40.506° from (002) of Cr2N was still present, indicating that the outer Cr2N layer was not totally removed after grinding in this case). It is possible that this phase change in the inner Cr-carbide layer took place at a coating time between 2 h and 4 h, which not only caused the two phenomena discussed above but

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Cr2N

Cr2N

Cr2N

5

Cr2N

Cr2N

Cr-carbide Cr-carbide

Cr-carbide Cr-carbide Cr-carbide

0h

0.5 h

10 µm

1h

10 µm

2h

4h

10 µm

10 µm

10 µm

Fig. 5. Microstructure of the coatings formed on steel 45# by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at 1100 °C for various lengths of time.

also the change in the slope of the thickness dT versus t1/2 plot as shown in Fig. 6a. It appears to be counterintuitive that both outer Cr2N and inner Crcarbide layers can form simultaneously. To find out how the formation sequence of these two layers is affected by processing temperature at the very initial stage of the coating formation process, a series of coating experiments were carried out to investigate the effect of processing temperature on the microstructures of the coatings formed by keeping the coating time constant at 0 h.

25

(a)

d T (µm)

20 y = 3.432x + 13.061

15 10

t = 2.7 h y = 7.5421x + 6.2639

5 0 0

0.5

1

1.5 t

1/2

2

2.5

4.3. Effect of processing temperature on the microstructures of the coatings The steel grade 45# was used again as the substrate. The processing temperatures investigated were 1000 °C, 1050 °C and 1100 °C. In all cases, the specimens were heated to the processing temperature using a heating rate of 10 °C/min and then immediately furnace-cooled at the natural cooling rate of the tube furnace. Thus, the coating time for all the specimens was 0 h. The microstructures of the coatings obtained are compared in Fig. 8. For specimens coated at 1050 °C and 1100 °C, apart from the difference in thickness of different layers, the microstructures were the same as those described in previous sections, i.e., consisting of an outer Cr2N layer and an inner Cr-carbide layer. For the specimen coated at 1000 °C, however, the continuous layer obtained was only the Cr-carbide layer and the Cr2N phase formed only in isolated areas on tope of the Cr-carbide layer. As coating time increases, a continuous Cr2N layer may also form eventually on the outmost surface. It is thus clear that both out Cr2N layer and inner Cr-carbide layer formed simultaneously at processing temperatures above 1050 °C, but when processing temperature was below this level, a continuous Crcarbide layer formed first on the surface, which may then be followed by the formation of a continuous Cr2N layer on top of the prior Cr-carbide layer as coating time increases. Thus, only when the processing temperature is low enough, can this layer growth sequence be observable. However, the pack cementation process is the one that requires thermal and kinetic activations. If the processing temperature is too low, the packs may not be adequately activated, and the diffusion and

3

1/2

(h ) 800

18

Cr2N(002)

(b)

Cr23C6

700

15

511

(844)

600 500 Intensity

d i (µm)

12 9

(422)

400

(420) 300

6 3

Cr2N

200

Cr-carbide

100

0 0

2

4

6

8

10

t (h) Fig. 6. Dependence of Cr2N and Cr-carbide layer thickness on coating time t; (a) plot of total Cr2N + Cr-carbide layer thickness versus t1/2, (b) plot of individual Cr2N and Crcarbide layer thickness versus t.

(533) (440) (600)

(820)

(622)

(800)

(751) (662)(931)

70

80

0 30

40

50

60

90

2θ (˚) Fig. 7. XRD pattern measured from a surface with the outmost Cr2N layer being removed by grinding for a coating formed on steel 45# by pack cementation in a pack 30Cr2N– 2NH4Cl–68Al2O3 (wt.%) at 1100 °C for 8 h (PDF#: 00-035-0783).

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(a)

Cr2N Cr-carbide

Cr2N

1000 ˚C

Cr2N

Cr2N

Cr-carbide

Cr-carbide

Cr-carbide

1050 ˚C

precipitates

1100 ˚C

Cr-enriched 5 µm

5 µm

5 µm

Fig. 8. Microstructure of the coatings formed on steel 45# by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at different temperatures by keeping the coating time constant at 0 h.

4.4. Effect of carbon content in steel substrate on microstructure of the coatings The steel grade Q235C was used as substrate for this investigation. The major difference in alloy composition between this steel grade and steel 45# used in previous sections is in the carbon content. The carbon content in the former is specified in the range of ≤0.17 wt.%, which is much lower than that in the latter, which is in the range of 0.42– 0.5 wt.%. The specimen was treated at 1100 °C for 4 h. The cross sectional SEM image and concentration depth profiles measured by EDS across the coating layer showed that the coating consisted of an outer Cr2N layer with a thickness of approximately 9 μm and a large Cr-enriched interdiffusion zone underneath; the thickness of this interdiffusion zone was about 37 μm (Fig. 9). The inner Cr-carbide layer as observed in the coatings formed on steel 45# as shown in previous sections was not present in this coating (Fig. 8a), and indeed in all the coatings formed on this steel grade. The phase of the outermost surface layer was also confirmed as Cr2N by XRD. It is thus clear that the carbon content in this steel grade was not high enough for forming a continuous Cr-carbide inner layer. Instead, the Cr-carbide was precipitated in the Cr-enriched interdiffusion zone in the form of nodular or needle-shaped particles (Fig. 8a). It is also evident that the surface showed a wavy profile and the thickness of the Cr-enriched interdiffusion zone was not uniform (Fig. 8a). These features were caused by the internal Cr-carbide precipitation in the Cr-enriched interdiffusion zone. As illustrated in Fig. 8a, the thickness of the Cr-enriched interdiffusion zone at the location marked out by the elliptic circle was noticeably smaller than at other locations. At this marked out location, a semi-continuous line of Cr-carbide precipitates was present across the thickness of Cr-enriched interdiffusion zone and it was oriented at a small angle to the thickness direction. The inward Cr diffusion would be slower at this location because of the presence of this semi-continuous line of Cr-carbide precipitates as Cr diffusion in Cr-carbide is likely to slower than in steel matrix. Consequently, the growth of the Cr-enriched interdiffusion zone will be slower at the locations where such a semi-continuous line of Cr-carbide precipitates is present than at locations where such a semi-continuous line of Cr-carbide precipitates was not present. As the growth of the Cr-enriched interdiffusion zone was accompanied by volume increase, the un-even inward Cr diffusion rates would lead to a non-uniform diffusion layer thickness across the surface plane, resulting in a wavy surface when viewed at the cross section under SEM. It is relevant to point out that, despite the shape change in the surface as discussed above, there was no spallation or microcracks in the outer Cr2N layer. On the one hand, this surface shape change took

100 Concentration (at.%)

reaction kinetics may become too low for the coating formation process to occur.

diffusion zone

30 µm

(b) 80 N

60

Cr 40

Fe

20 0 0

10

20

30

40

50

Depth (µm) Fig. 9. Cross-sectional SEM and element depth profiles measured by EDS for a coating formed on steel Q235C by pack cementation in a pack 30Cr2N–2NH4Cl–68Al2O3 (wt.%) at 1100 °C for 4 h.

place at processing temperatures at which the steel substrate would become quite soft and hence stresses raised in the coating as different coating layers grew would be relaxed by creep in steel substrate at high temperatures. On the other hand, the Cr concentration showed a decreasing depth profile in the Cr-enriched interdiffusion zone (Fig. 8b), which provided an intermediate layer in which thermal expansion changed gradually in the thickness direction, which alleviated the adverse effect of residual thermal stresses arising during cooling and hence ensured the integrity of the outer Cr2N layer. In fact, these mechanisms operated in all the coatings described in previous sections. As described previously, the outer Cr2N layer thickness on steel Q235C after coating at 1100 °C for 4 h was approximately 9 μm. This was much thinner than that formed on steel 45# (about 17 μm) under the same processing conditions (see Section 4.1). Indeed, it was observed in general in this study that the outer Cr2N layer grew more easily on steel 45# than on steel Q235C, despite the fact that an inner Crcarbide layer always formed in the surface of the former steel grade, but not in the surface of the latter steel grade. Thus, the formation of the inner Cr-carbide layer did not hinder but facilitated the growth of the outer Cr2N layer. 5. Conclusions The formation of chromium nitride coating on carbon steels by pack cementation process has been studied as a function of processing time, temperature and carbon content in steel substrate. On the basis of the results obtained, the following conclusions may be made: 1. It is feasible to form a continuous layer of chromium nitride on carbon steels by pack cementation process in the temperature range 1000–1100 °C using NH4Cl as an activator, Cr2N nitride powder as a

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source for both nitrogen and chromium and Al2O3 power as inert filler. If the carbon content in steel is high enough, the coatings formed would have a microstructure consisting of an outer Cr2N layer, an inner Cr-carbide layer and a Cr-enriched interdiffusion zone underneath the inner Cr-carbide layer; in this case, both outer Cr2N layer and inner Cr-carbide layer would form and grow simultaneously at temperatures higher than 1050 °C, although it is possible that the inner Cr-carbide layer would form first during the heating stage. The phase of the inner Cr-carbide layer may change with coating time from a carbide with a lower Cr concentration to a carbide with a higher Cr concentration. This change in Cr-carbide phase in the inner layer would alter the growth kinetics of the coating, particularly the kinetics of the inner Cr-carbide layer. If the carbon content in steel substrate is not sufficiently high, the inner Cr-carbide layer may not form and the microstructure of the coating would only consist of an outer Cr2N layer with a Cr-enriched interdiffusion zone underneath in which Cr-carbide precipitates in the form of nodular and needle-shaped particles. The outer Cr2N layer maintained its integrity without any spallations and microcracks occurring during cooling due to the fact that the Crenrich interdiffusion zone provided an intermediate layer in which the gradual change in alloy composition in the thickness direction offered a thermal expansion gradient that reduced the harmful residual thermal stress effects. Growth of the outer Cr2N layer would be faster on a steel substrate in which an inner Cr-carbide layer can simultaneously form than on a steel substrate in which such an inner Cr-carbide layer cannot form.

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Please cite this article as: X.J. Lu, Z.D. Xiang, Surf. Coat. Technol. (2016), http://dx.doi.org/10.1016/j.surfcoat.2016.10.047